Jan 15, 1991 - Ienergy input based upon their electrical resistivities, see .... of processes in which electrical energy is dissipated within the powder mass to cause .... The basic principles of these rotating electrical machines as pulsed power ...... J. M. Chou, "Electrical Resistance Sintering of Motor/Generator Brushes Under.
I U CONTROLLING FUNDAMENTALS IN HIGH-ENERGY HIGH-RATE PULSED POWER MATERIALS PROCESSING OF POWDERED TUNGSTEN, TITANIUM ALUMINIDES, AND COPPER-GRAPHITE COMPOSITES
3
3
IN
Principal Investigators: Harris L. Marcus and William F. Weldon
*
Co-Investigators: Chadee Persad, Zwy Eliezer and David Bourell
Center for Materials Science and Engineering & Center for Electromechanics
I' JAN
U
!E
I
15 1991
The University of Texas at Austin Austin, Texas
I
October 1990 UTCMSE-90-02 Final Technical Repo
r-
i'
A
Army Research Office Contract DAAL 0387-K-0073
44XDVed for public relece;
0
tr
I I I
I
Final Technical Report Contract DAAL 0387-K-0073 The University of Texas at Austin Reporting Period: June 1, 1987 - August 31, 1990
I I I I CONTROLLING FUNDAMENTALS IN HIGH-ENERGY HIGH-RATE PULSED POWER MATERIALS PROCESSING OF POWDERED TUNGSTEN,
I I I
TITANIUM ALUMINIDES, AND COPPER-GRAPHITE COMPOSITES
Principal Investigators
I I
Harris L. Marcus
William F. Weldon
Professor of Mechanical Engineering Director of Center for
Professor of Mechanical & Electrical Engineering Director of Center for Electromechanics
Materials Science and Engineering Report Number: UTCMSE-90-02
I *o
IPrv~d
puji
=
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I I1TITLE (include Securify Class CONTROLLJNG FUNDAMENTALS INHIGH-ENERGY HIGH-RATE PULSED POWER MATERIALS PROCESSING OF POWDERED TUNGSTEN, TITANIUM ALUMNDES, AND COPPER-GRAPHITE COMPOSITES [FINAL REPORT. UNCLASSIFIED]J 12. PERSONAL AUTHOR(S)
C. PERSAD, H. L. MARCUS, D. L. BOURELL, Z. ELIEZER. W. F.WEDN 114 DATE OF REPORT (erMot.Dy
13.TIECVERED FRM81UN 01
13a. TYPE OF REPORT
jCOUNT
90_OCT31_P1_ TO90 AUG 31 REPORT *FINAL * 16 SUPPLEMENTARY NOTATION The view, opinions and/or findings contained in this report are those of the auth r(s).and sh uld not ,be~const u d as an ff1Aiial De armnt of the Army position, 18. SUBJECT TERMS (Continue on reverse if necessary and identify by block number)
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19 ABSTRACT (Conltinlue on reverse it necessary and identify by block number)
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18.
SUBJECT TERMS (BASED UPON MATERIALS RESEARCH SOCIETY KEYWORDS BY TOPIC GUIDELINES)
A.
SYNTHESIS: chemical synthesis, powder metallurgy, self-propagating high-temperature synthesis, sintering, sol-gel.
B.
PROCESSING: bonding, extrusion, forging, welding.
C.
MATERIAL TYPE: alloys-intermetallic, ceramics, composites, glasses-metallic, metals, superconductors.
D.
MATERIAL FORM: colloidal, laminate, mixture, polycrystal, powder.
E.
ANALYSIS TECHNIQUES: Auger electron spectroscopy, computer simulation, differential thermal analysis, hardness testing, optical metallography, scanning electron
microscopy, scanning/transmission electron microscopy, transmission electron microscopy, x-ray diffraction. F.
PROPERTY STUDIED: electrical properties, mechanical properties, tribology, crystallographic structure, phase transformation, diffusion.
19. ABSTRACT
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X'
This study was conducted to determine the controlling fundamentals in the high-energy high-rate (1 MJ in Is) processing of metal powders. This processing utilizes a large electrical current pulse to heat a pressurized powder mass.The current pulse was provided by a homopolar generator. Simple short cylindrical shapes were consolidated so as to minimize tooling costs. Powders were subjected to current densities of 5 kA/cm2 to 25 kA/cm2 under applied pressures ranging from 70 MPa to 500 MPa. Disks with diameters of 25 mm to 70 mm, and thicknesses of 1 mm to 10 mm were consolidated. Densities of 75% to 99% of theoretical values were obtained in powder consolidates of tungsten, titanium aluminides, copper-graphite, and other metal-ceramic composites. Extensive microstructural characterization was performed to follow the changes occuring in the shape and microstructure of the various powders. The processing science has at its foundation the control of the duration of elevated temperature exposure during powder consolidation The goal was a fuller understanding of the interrelated electrical, thermal and mechanibal interactions that occur during the observed rapid densification.Three major densification mechanisms were observed. These were: plastic flow at elevated temperature in singlephase systems such as tungsten; plastic flow of a minor low-temperature phase in a two phase non-interacting system such as tungsten-copper; and liquid-phase-assisted densification via the production of a small quantity of liquid within the consolidated mass. The process appears particularly well-matched to the elevated temperature, solid state consolidation of powders of metals, and of metal-matrix composites such as copper + graphite, in which exothermic reactions are absent. With the continued maturation of the now commercially available homopolar generator pulsed power sources, the materials processing technology flowing from this research will find specialized niches in the production of components that fully exploit its potential for single-residence conversion of powders to finished parts. High performance sliding electrical contacts provide one example of such a process-product match in which technology transfer is now imminent.
.
UN('LASSIF1ED SECUR]ITY CLASSIFICATION OF THIS PACE
I I 1. FOREWORD
This is the final report of work completed under DARPA/ARO Contract DAAL 0387-K-0073, [June 1, 1987 - August 31, 1990], to perform research on the controlling fundamentals of High-energy igh-rate materials processing using pulsed power sources. The work was performed by the Center for Materials Science and Engineering, and the Center for Electromechanics, Bureau of Engineering Research, The University of Texas at Austin, Austin, TX 78712.
Special materials were provided for this research by outside laboratories. Tungsten powders were received from Los Alamos National Laboratories, NM. (H. Sheinberg), and
i
from GTE. Titanium aluminide powders (7-TiAl and a2-Ti3Al-based) were provided by Ralph Anderson's group at Pratt and Whitney, FL.
Project monitors were E. Farnum, DARPA, and A. Crowson, ARO. Accession For
NTIS GRA&I
DTIC TAB Unannounced Justification
Q
,Distribution/
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I
Availability Codes Avail
IDist
and/or
Special
2. TABLE OF CONTENTS
I
1.
Foreword
2.
Table of Contents
ii
3.
List of Tables, Figures, and Appendices
iii
4.
MAIN REPORT
3
I
I
I
Page
4A.
Statement of Problem
1
4B.
Summary of Results
2
4.1
Comparison with Other Processes
2
4.2
Historical Background
7
4.3
Experimental Apparatus
8
4.4
Densification Mechanisms
11
4.6
Towards a Manufacturing Technology
12
4.7
Specific Materials Systems
13
4C.
List of Publications and Technical Reports
17
4D.
List of Scientific Personnel
22
5.
Bibliography
6.
Appendices
24
6A
Processing Overviews
25
6B
Tungsten Processing
42
6C
Titanium Aluminide Processing
116
6D
Copper-Graphite Processing
141
6E
Metallic Glasses
203
6F
Electrotribology
232
6G
Other Systems
265
ii
I I 3. LIST OF TABLES, FIGURES, AND APPENDICES
I
U
I I
Page
Table 4.1
A comparison of pressure-assisted densification processes.
3
Table 4.2
Materials Systems Studied in this Research and Published Sources.
13
Figure 1
A power density versus interaction time plot indicating the fields of IEHR processing research. PPP1 and PPP2 indicate the fields in which this HEHR processing research was performed. PPP3 is the regime for pulsed power sources 5
now under development.
I
3
Figure 2
A
IT diagram demonstrates the potential advantage of HEHR
consolidation in maintaining non-equilibrium structures. A profile for a two-phase powder assemblage is shown thermally pulsed
3 I
U 3
s) high temperature exposure. with a short (< 1
6
Figure 3
Schematic of HEHR consolidation processing apparatus.
9
Figure 4
Power spectra of the (a) as-received powder consolidates, (b) low (2000 J/g) input energy samples #757 (AR), #765 (1), #827 (II). Metallic glass powders in various stages of crystallization show sensitive responses to
3 I
Ienergy
input based upon their electrical resistivities, see
Appendix E-3 for a detailed discussion. iii
10
I Appendices Appendix A: Processing Overviews A-1.
"25
High-Energy High-Rate Materials Processing
.26
A-2. Single Residency Sintering and Consolidation of Powder
Metal Alloys, Intermetallics, and Composites by Pulsed Homopolar Generator Discharge
Appendix B: Tungsten Processing
3
31
42
B-1. Manufacturing Metallurgy of Tungsten: A Novel Integrated Powder Metallurgy Approach
43
B-2. Development and Control of Microstructure in P/M Tungsten by Synthesis and Reduction of Tungsten Trioxide Gels
I
50
B-3. High-energy, High-rate Consolidation of Tungsten and Tungsten-based Composite Powders
58
B-4. Synthesis and High Energy High Rate Processing of Ultrafine Grained Tungsten
69
Appendix C: Titanium Aluminide Processing
116
C-1. High-Energy High-Rate Processing of High-Temperature
I
Metal-Matrix Composites
117
C-2. Microstructure/Processing Relationaships in High-Energy
3
High-Rate Consolidated Powder Composites of Ti3A1 + TiAl
I
iv
123
II i
I
C-3. Matrix-Reinforcement Interface Characteristics of (Ti3AI + Nb)-Based Powder Composites, Consolidated by High-Energy High-Rate Processing
I
Appendix D: Copper-Graphite Processing
I
3
D-2. A Novel Processing Technique for Metal/Ceramic Composites
I
D-3. Advanced Composite Materials for High-Performance Electrotribological Applications
Appendix E: Metallic Glasses E-1.
167
172
203
204
E-2. Identification of an eta-Boride Phase as a Crystallization
I
Product of a NiMoFeB Amorphous Alloy
209
E-3. High Energy-High Rate Powder Processing of a Rapidly-
I
Quenched Quaternary Alloy, Ni 5 6 .5 Mo 23 .5 Fe 10 B 10
Appendix F: Electrotribology F-1.
Iv
142
Consolidation of Metallic Glass Ribbons using Electric Discharge Welding
3
141
D-1. High-Energy/High-Rate Consolidation of Copper-Graphite Composite Brushes for High-Speed, High-Current Applications
I
129
Composite Solid Armatures for Railguns
215
232 233
I I F-2. Composite Solid Armature Consolidation by Pulse Power Processing: A Novel Homopolar Generator Application in EML Technology
F-3. Wear of Conductors in Railguns: Metallurgical Aspects
I
247
251
F-4. Nanosized Structures in Cu-W-WC-C Composites for Electrotribological Applicaz!ins
Appendix G: Other Systems
256
265
G-1. Basic Principles for Selecting Phases for High Temperature Metal Matrix Composites: Interfacial Considerations
I
3
266
G-2. Consolidation of Powders of the Oxide Superconductor YBa2Cu3Ox by High-Energy High-Rate Processing
270
G-3. High-Energy High-Rate Deformation Processing of
I
P/M Aluminum-Silicon Carbide Composites
286
G-4. Electrical Characteristics of High Energy- High Rate
I
Consolidated Metal Matrix Composites
G-5. High Rate Consolidation of Co-Fe-B Ferromagnetic Powders
I I I
vi
291
315
I
I 4A. STATEMENT OF THE PROBLEM STUDIED
The objective of this three-year effort was to develop an understanding of the controlling fundamentals in high-energy high-rate materials processing using pulsed power. The primary powder materials systems selected for study were tungsten, titanium aluminide, and copper-graphite. Other systems were investigated to define specific niches
I
I I I I I I I I I I I I
where this type of processing could be advantageously applied.
I1
I I I
4B.
SUMMARY OF THE MOST IMPORTANT RESULTS
INTRODUCTION
This final report is a summary and compilation of published research results. The research was funded by DARPA and monitored by ARO. The research was aimed at
I
understanding the con
lling fundamentals in high-energy high-rate powder processing.
In this process, pressurized conductive powders are densified by the Joule (12R) heating obtained from a current pulse. The large current pulse is provided by the discharge of a Homopolar Generator (HPG). This section summarizes the findings of the research into the controlling
I
fundamentals in High-Energy High-Rate Processing. In section 4.1, the processing method is compared to the global set of powder consolidation processes. In section 4.2, its historical development is traced. In section 4.3, the methodology associated with the application of the technique in laboratory research is described. In section 4.4, the observed mechanisms that control densification are discussed. Factors pertinent to the evolution of
I
this process into a manufacturing technology are indicated in section 4.6. In section 4.7, results obtained on specific materials systems are described.
I I I I I
I 4.1 Comparison with Other Processes
Efficient densification of metallic powders can be achieved by the simultaneous application of temperature and pressure. Uniaxial vacuum hot pressing, hot isostatic pressing, and hot large-strain deformation processing such as powder rolling and powder extrusion which exploit this approach are well-established processes for making fully dense
I
I
materials from powder.
High-energy High-rate processing for powder consolidation belongs to the group of processes in which electrical energy is dissipated within the powder mass to cause a rise in temperature. When combined with pressure application, this process leads to rapid
I
densification.
In powder consolidation, the pressure-assisted densification processes have been differentiated by the magnitude, duration, and directionality of the applied pressure.
I
Table 4.1. A comparison of pressure-assisted densification processes [I].
PROCESS
MAGNITUDE
DURATION (s)
(GPa)
I
I3
DIRECTIONALITY OF LOADING
Hot pressing
0.01-0.03
103 -104
Uniaxial
Hot Isostatic Pressing
0.10-0.30
103 -104
Isostatic
Hot Extrusion
0.10-1.00
102- 1(4
Complex
HEHR
0.10-0.50
1-10
Uniaxial
Explosive
10-100
- 10.6
Complex
Table 4.1 provides a comparison of several of the mature powder consolidation processes with the High-Energy High-Rate (HEHR) Process. It is observed that a key feature of the HEHR process is the short duration of the time at elevated temperature. By careful matching of a pressure application profile to the temperature developed within the powder mass, efficient pressure-assisted densification can be obtained by plastic flow. Appendix A-I provides an overview of the metallurgical approach to the application of pulsed power to materials processing. The interaction time and power density mapping of a variety of materials processes are mapped in Figure 1. Overlaid is the regime labelled PPP1 in which most of this research was performed. PPP2 is the regime for capacitorbank powered consolidation. Less conductive materials such as the YBCO superconductor powders were processed in this regime [see Appendix G-2 for details].
I
From a metallurgical standpoint, the manipulation of time-at-temperature is the common approach to control of transformations. Figure 2 indicates the potency of HEHR processing in controlling transformations by limiting the time at elevated temperature. Appendix A-2 discusses some of the engineering aspects associated with implementation of this type of processing.
I I I I I I I4
I
Low Inductance
Compulsators
Homopolar
Pulse Forming Networks
i
100 8 10 -8
"6
P
Generators
-4
10 0
1
10 .
10
9
10
1
10
9°
SPECIFIC ENERGY J/cm PPP2
10
-
10
SHOCK HARDENING
E
8
8
10
10 3:
I
H 17 >d
RAPID
-
SURFACE ALLOYING AND
Fn
SMELZI6 1 . ?-0
I4
0
10
DILLING PP10 PPP3 6
-10
10 WELDING AND CUTTING
5
0 10 a.
10
5
.CLADDING
AND
10 10
i
10
? -PPP
-
4
HARDENING 3
o10 10*
: 10.6
10.4
INTERACTION TIME
I
-
10-2 SECONDS
3
1
10 10
Figure 1.A power density versus interaction time plot indicating the fields of HEHR processing research. PPPI and PPP2 indicate the fields in which this HEHR processing research was performed. PPP3 is the regime for pulsed power sources now under development.
5
I I I I I I
LIQUID
*
TMB
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I"EQUILIBRIUM" ISTRUCTURE
j "DESIRABLE" STRTICTIIRE
HIGH-ENERGY HIGH-RATE PROCESSING REGIME
I
< 1 SEC.
TIME
I I I
Figure 2. A IT diagram demonstrates the potential advantage of HEHR consolidation in maintaining non-equilibrium structures. A profile for a two-phase powder assemblage is shown thermally pulsed with a short (< 1 s) high temperature
*
*
exposure.
I
I
4.2 Historical Backgound
Powder-based metallic materials and metal-ceramic composites can be produced by a variety of methods including the use of large electrical currents. For over fifty years it has been known that the direct electrical resistance heating of conductive metallic powder materials can cause densification. The mechanism by which this densification proceeds is material specific and is related to time at temperature during processing. The interest in the use of electrical current for heating of powders has been sporadic [2-5] and has been
I
I
handicapped by the absence of a well-matched power source.
Over the last decade, devices for the production of pulsed electrical power have enjoyed a developmental thrust. In particular, homopolar generators (HPG) based upon the Faraday disk have been successfully engineered and commercialized [6]. The availability of these machines as pulsed power sources has fostered the development of novel powder processing approaches for metals and metal-ceramic composites.
At the University of Texas at Austin, powder processing powered by a homopolar generator has been developed as a high-energy high-rate (HEHR) materials processing technique [7]. The '1 MJ in Is' pulsed energy delivery from a 10 MJ HPG has been employed in a wide-ranging series of processing experiments, including welding, billet
I
heating, and powder consolidation.
I
I
4.3 Experimental Apparatus
In laboratory research, the two key components in the apparatus are a powde heating source and a powder pressurization source. These are shown schematically in Figure 3 in conjunction with the die. The powder mass is maintained under pressure in the die between two electrodes which also deliver the electrical current for heating.
I
As indicated earlier the powder heating occurs by pulsed Joule Heating [proportional to the product of the square of the current and the resistance of the powder mass (12R)]. The heating source is therefore a high current delivery system - a Homopolar generator. The basic principles of these rotating electrical machines as pulsed power supplies have been discussed and compared to other pulsed power systems by Weldon [8]. For powder consolidation, these machines are capable of delivering high stored energies via currents at 106 Ampere levels. Their low voltage (< IOOV) is sufficient for the processing of most metallic-conducting systems.
The pressure is provided by a hydraulic press with 100 ton capacity. By incorporating customized accumulators and timers into a basic system, it was possible to increase the pressure at a time during the temperature cycle corresponding to the peak temperature, thus obtaining plastic-flow induced densification at minimum flow stress. While the system employed in this research was based upon preset values of pressure vs time, a fast-response system with feed-back based upon changes in resistivity of the powder mass may be easily engineered into a system for a specific industrial application.
I
An example of the clear changes in resistivity during the processing of metallic glass powders is presented in Figure 4 and discussed in Appendix E-3.
*
8
I I I
U
Fixed platen
Electrode
HPG
Powdr
I Backup ring Hydraulic press
i I
i i
Figure 3. Schematic of HEHR consolidation processing apparatus.
I I I I I i
9
200
I
3
I
1
-
2
IO0
IWi
ENERGY
721 820
1850 jig 1970
752
2285
a
00. .0
0
m
SPECIMEN
~so'%
0-
0.5
0.0
1.0
1.5
PULSE DURATION (sec)
b
200
/
150
AR
WU100
0 _
0
0.0
0.5
1.0
1.5
PULSE DURATION (sec)
I CLC
50
00
0.5
1.0
1.5
PULSE DURATION (sec)
I
Figure 4. Power spectra of the (a) as-received powder consolidates, (b) low (2000 J/g) input energy samples #757 (AR), #765 (I), #827 (II). Metallic glass powders in various stages of crystallization show sensitive responses to energy input based upon their electrical
3
resistivities, see Appendix E-3 for a detailed discussion.
3
10
4.4 Densification Mechanisms
Different material-specific mechanisms were observed to operate in the singleresidence conversion of loose powders into a consolidated mass of near full density. Three broad classes are distinguished: (i) single-phase systems processed in the solid state, eg. W, (ii) two-phase systems in which deformation occurs preferentially in the lower strength
I
phase, eg., W-Cu, (iii) multiphase systems in which a transient liquid phase is formed during processing either by a solid-4liquid phase change of one constituent, or by the formation of a reaction product with lower melting temperature than the parent phases, eg., W-Ni, W-Ni-Fe + B4 C. Most of this research was conducted with systems in which the materials were maintained in the solid state, and in which the densification occured by plastic-flow. This type of HEHR processing is essentially a single-residence closed-die powder forging operation in which the elevated exposure temperature is limited to very short times.
I
From a materials synthesis standpoint, the processing of multi-phase systems in which new liquid phases are produced offers many exciting possibilities. From an engineering standpoint however, the containment of liquids at high pressures requires considerably more effort in die design and in the selection of die materials that resist attack by liquid metals. One interesting approach to this problem is to utilize the liquid formed in an exfiltration mode. In this miethod a controlled volume of liquid is formed well within the
3
hot pressurized mass and forced outward in all directions becoming solid again before it
I
can exit from the consolidated mass.
I
11
i
I
I
4.6 Towards a Manufacturing Technology
HEHR powder processing is still a research laboratory technique awaiting industrial deployment. The vehicle for such deployment is expected to be a product which utilizes the technique at its full potential. Two such products are now under consideration, both of which have origins in this research. These are a high-speed, high-current sliding electrical
I
contact based on copper-graphite, and a wear-resistant mining tool based upon cem.1nted tungsten carbide. Both of these attempt to exploit a unique feature of HEHR processing: the ability to consolidate powders and to bond the consolidate onto a wrought mandrel all in the same process.
I I I I I I I I I I *
12
I
I
4.7 Specific Materials Systems
The focus of the program has been an experimental effort to define processing/ microstructure/property relationships. Although the emphasis was on understanding the processing science, several materials systems of varying complexity were studied.
I
Three major materials systems were studied: tungsten; titanium aluminides; and copper-graphite composites. The highlights of the results are given here. Details are provided in the published papers which are appended. Table 4.2 is a compilation of the many materials systems that were investigated and the sources in which the results are
I
documented.
Table 4.2. Materials Systems Studied in this Research and Published Sources
I
I
*13
PUBLISHED SOURCE
MATERIALS SYSTEM
Al-SiC
Appendix G-3, Appendix G-4
TiAl Ti3 Ai
Appendix C-2, Appendix C-3 Appendix C-i, Appendix C-2, Appendix C-3
Fe-Si-B
Appendix E- 1
Co-Fe-B Ni-Mo-Fe-B
Appendix G-5 Appendix E-2, Appendix E-3
Cu-C
Appendix D-1, Appendix D-2, Appendix D-3
Cu-W Cu-W-WC-C YBa 2 Cu3O7 .x
Appendix F-2 Appendix F-4 Appendix G-2
W
Appendix B-1, Appendix B-2, Appendix D-3, Appendix B-4
W-Ni
Appendix B-3
W-Cu W-Nb W-Ni-Fe + B4 C
Appendix D-3 Appendix B-3 Appendix B-i, Appendix B-2, Appendix B-3
I'
In the case of the tungsten processing, a fully integrated approach was developed that included the synthesis of ultrafine metallic powders from oxide precursors. These ultrafine powders provided excellent matching to the process since the resistance of the powder mass constituted therefrom had a large constriction resistance contribution. This particle-size controlled component of the resistance improved the efficiency with which the
I
powder mass could be heated, and allowed the achievement of the high temperatures necessary for tungsten consolidation. Results are detailed in Appendix B-I to B-4.
The initial effort on the titanium aluminides examined the processability of y-TiAl powders, and of Nb-stabilized Ti 3AI. Because of their inherent low ductility, the y-TiAl disks produced were always cracked, often in a radial pattern. Crack-free disks of Nbstabilized Ti3Al were produced at near-full density. The latter composition was then employed as an intermetallic matrix and reinforced with ceramic powders such as SiC, TiB 2 and AIN. The interfaces produced in these systems showed a variety of reaction products indicating that these reinforcements were not stable at high temperatures. An
I
effort to produce a composite of TiAl dispersed in Ti3Al matrix showed a more stable interface. Details are given in Appendix C-1 to C-3.
The effort on copper-graphite composites was aimed at the consolidation of fully dense material from elemental copper and carbon powders. It was anticipated that such
I
binderless consolidation would allow these materials, as sliding electrical contacts, to surpass the 200m/s barrier that had been observed in conventionally processed materials which contained Pb-Sn additives. This performance objective was met by the materials consolidated by this method [9]. A continuing program funded by the State of Texas is now underway to commercialize this class of miaterials for high performance sliding
I
electrical contact applications. 14
I !4 Other significant efferts were in the areas of metallic glass processing [see Appendix E] and in the area of solid armatures/electrotribology [see Appendix F]. The effort on metallic glasses led to the discovery of a new eta-boride phase, (see Appendix E-2), and the understanding of the influence of degree of crystallinity upon the heat evolved during processing and its effect on the microstructure of HEHR consolidated material (Appendix E-3). The solid armatures/electrotribology area provided interesting insights into how materials respond to large pulsed currents in electromagnetic launch technology and contributed to the understanding of the possibilities in multi-phase materials processing using pulsed power. These insights have also produced some new initiatives aimed at the development of specially tailored materials for these high performance electrotribological applications (for example, see Appendix F-4).
II I I I I I I I *
15
Summay
I
This study was conducted to determine the controlling fundamentals in the highenergy high-rate (1 MJ in Is) processing of metal powders. This processing utilizes a large electrical current pulse to heat a pressurized powder mass. The current pulse wis provided bv a homopolar generator. Simple short cylindrical shapes were consolidated so as to minimize tooling costs. Powders were subjected to current densities of 5 kA/cm2 to 25
I
kA/cm2 under applied pressures ranging from 70 MPa to 500 MPa. Disks with diameters of 25 mm to 70 mm, and thicknesses of 1 mm to 10 mm were consolidated. Densities of
3
75% to 99% of theoretical values were obtained in powders of tungsten, titanium aluminides, copper-graphite, and other metal-ceramic composites. Extensive microstructural characterization was performed to follow the changes occuring in the shape
3and
microstructure of the various powders. The processing science has at its foundation the
control of the duration of elevated temperature exposure during powder consolidation. The
3
goal was a fuller understanding of the interrelated electrical, thermal and mechanical interactions that occur during the observed rapid densification.Three major densification mechanisms were observed. These were: plastic flow at elevated temperature in singlephase systems such as tungsten; plastic flow of a minor low-temperature phase in a two phase non-interacting system such as tungsten-copper, and liquid-phase-assisted
3 3
densification via the production of a small quantity of liquid within the consolidated mass. The process appears particularly well-matched to the elevated temperature, solid state consolidation of powders of metals, and of metal-matrix composites such as copper + graphite, in which exothermic reactions are absent. With the continued maturation of the now commercially available homopolar generator pulsed power sources, the materials
3 3
processing technology flowing from this research will find specialized niches in the production of components that fully exploit its potential for single-residence conversion of powders to finished parts. High performance sliding electrical contacts provide one example of such a process-product match in which technology transfer is now imminent. 16
4C. LIST OF PUBLICATIONS AND TECHNICAL REPORTS
1.
H. L. Marcus, D. L. Bourell, Z. Eliezer, C. Persad and W. Weldon, "High-Energy High-Rate Materials Processing," Journal of Metals, Vol. 39, No. 12, Dec. 1987,pp. 6-10.
2.
3
G. Elkabir, C. Persad, H. L. Marcus, and T. A. Aanstoos, "High-Energy High-Rate Consolidation of Aluminum- Silicon Carbide Composites", Proceedings of the Ninth Annual Discontinuously Reinforced MMC Working Group Meeting, Park City,
3
UTAH, January 1987, (MMCIAC No. 000695), MMCIAC, Santa Barbara, CA, 1988, pp. 363-386.
3.
Armatures for Railguns", J. H. Gully (ed.), Paper #13 in Proceedings of the 3rd
3 I
C. J. Lund, C. Persad, Z. Eliezer, D. Peterson, and J. Hahne, "Composite Solid
International Conference on Current Collectors, Austin, TX, November 1987.
4.
M. J. Wang, C. Persad, Z. Eliezer, and W. F. Weldon, "High-Energy/High-Rate Consolidation of Copper-Graphite Composite Brushes for High-Speed, HighCurrent Applications", J. H. Gully (ed.), Paper #20 in Proceedings of the 3rd
3 I
International Conference on Current Collectors, Austin, TX, November 1987.
5.
Y. W. Kim, D. L. Bourell, and C. Persad, "Consolidation of Metallic Glass Ribbons using Electric Discharge Welding", Metall. Trans., Vol. 19A, June 1988, pp 1634-1638.
I *
17
6.
M. E. Fine, D. L. Bourell, Z. Eliezer, C. Persad, and H. L. Marcus, "Basic Principles for Selecting Phases for High Temperature Metal Matrix Composites: Interfacial Considerations" Scripta Met., 22, June 1988, pp. 907-910.
7.
Y. W. Kim, L. Rabenberg, and D. L. Bourell, "Identification of an eta-Boride Phase as a Crystallization Product of a NiMoFeB Amorphous Alloy", Journal Mater. Res.
I
3(6), November/December 1988, pp. 1336-1341.
8.
3H.
C. Persad, S. J. Lee, D. R. Peterson, J. S. Swinnea, M. Schmerling, K. M. Ralls, L. Marcus and H. Steinfink, "Consolidation of Powders of the Oxide Superconductor YBa2Cu3Ox by High-Energy High-Rate Processing", Proceedings
Iof
the World Congress on Superconductivity, Houston, TX, February, 1988, C. G. Burnham & R. Kane (eds), Vol. VIII of World Scientific Series on PROGRESS IN HIGH TEMPERATURE SUPERCONDUCTIVITY, World Scientific Publishers, Teaneck, NJ, 1988, pp 460-475.
9.
Z. Eliez .r, M. J. Wang, C. Persad, and J. H. Gully, "A Novel Processing Technique for Metal/Ceramic Composites", in Ceramic Developments, C. C. Sorrell and B. Ben-Nissan (eds.), Materials Science Forum, Vol. 34-36, Trans. Tech.
3
3
Publ., Switzerland, (1988) pp 505-509.
10.
C. Persad, S. Raghunathan, B. H. Lee, D. L. Bourell, Z. Eliezer and H. L. Marcus, "High-Energy High-Rate Processing of High-Temperature Metal-Matrix Composites", Proceedings of the MRS Symposium on High Performance
3
Composites, Reno, NV, April 1988, published as MRS Symposia Proc. Vol. 120, F. D. Lemekey, A. G. Evans, S. Fishman, and J. R. Strife (eds.), MRS,
3
3
Pittsburgh, PA, 1988, pp 23-28.
18
I
11.
C. Persad, D. R. Peterson, and R. C. Zowarka, "Composite Solid Armature Consolidation by Pulse Power Processing: A Novel Homopolar Generator Application in EML Technology", IEEE Trans. on MAgnetics, January 1989, pp 429-432.
I
12.
C. Persad, C. J. Lund, Z. Eliezer, D. Peterson, J. Hahne, and R. C. Zowarka, "Wear of Conductors in Railguns: Metallurgical Aspects", IEEE Trans. on Magnetics, January 1989, pp 433-437.
I 13.
I
H. L. Marcus, C. Persad, Z. Eliezer, D. L. Bourell, and W. F. Weldon, "HighEnergy High-Rate Processing of Composites", Proceedings of the Tenth Annual Discontinously Reinforced MMC Working Group Meeting, Park City, Utah, January 1988, (MMCIAC No. 000696), MMCIAC, Santa Barbara, CA, 1989, pp 235-259.
I
14.
C. Persad, S. Raghunathan, A. Manthiram, M. Schmerling, D. L. Bourell, Z. Eliezer, H. L. Marcus, and W. F. Weldon, "Manufacturing Metallurgy of Tungsten: A Novel Integrated Powder Metallurgy Approach", IIE/OAS/SwRI Proceedings of
3
the Tenth Inter-American Conference on Materials Technology, April, 1989, San
I
Antonio, TX, Volume II, Section 30, pp 33-39.
15.
3
H. M. Tello, C. Persad, H. L. Marcus and D. L. Bourell, "High-Energy High-Rate Deformation Processing of Aluminum-Silicon Carbide Composites", IIE/OAS/SwRI Proceedings of the Tenth Inter-American Conference on Materials Technology, April, 1989, San Antonio, TX, Volume II, Section 30, pp 27-3 1.
I I
19
I 16.
I
W. F. Weldon and T. A. Aanstoos, "Single Residency Sintering and Consolidation of Powder Metal Alloys, Intermetallics, and Composites by Pulsed Homopolar Generator Discharge", Journal of Mechanical Working Technology, 2Q (1989) 353-
I
363.
17.
I
C. Persad, B. H. Lee, C. J. Hou, Z. Eliezer, and H. L. Marcus, "Microstructure/ Processing Relationaships in High-Energy High-Rate Consolidated Powder Composites of Ti3Al + TiAl", MRS Symposium Proceedings, Vol. 133: "High Temperature Ordered Intermetallic Alloys III", edited by C. T. Liu, A. I. Taub, N. S. Stoloff, and C. C. Koch, MRS, Pittsburgh, PA, 1989, pp 717-722.
I
18.
Y. W. Kim, D. L. Bourell and C. Persad, "High Energy-High Rate Powder Processing of a Rapidly-Quenched Quaternary Alloy, Ni 5 6 .5 Mo 2 3 .5 Fe 10 B10 ", Materials Science and Engineering, A 123, 1990, pp. 99-115.
I 3
19.
C. Persad, S. Raghunathan, A. Manthiram, M. Schmerling, D. L. Bourell, Z. Eliezer, and H. L. Marcus, "Development and Control of Microstructure in P/M Tungsten by Synthesis and Reduction of Tungsten Trioxide Gels", Solid State Powder Processing, A. H. Clauer and J. J. de Barbadillo, eds., TMS, Warrendale,
3
PA, 1990, pp. 357-364.
20.
Z. Eliezer, B.-H. Lee, C. J. Hou, C. Persad, and H. L. Marcus, "MatrixReinforcement Interface Characteristics of (Ti3 AI + Nb)-Based Powder Composites, Consolidated by High-Energy High-Rate Processing", in Metal & Ceramic Matrix Composites: Processing. Modeling & Mechanical Behavior, R. B. Bhagat, A. H. Clauer, P. Kumar, and A. M. Ritter, eds., TMS, Warrendale, PA, 1990, pp. 401412.
20
I I I
21.
G. Elkabir, C. Persad and H. L. Marcus,"Electrical Characteristics of High EnergyHigh Rate Consolidated Metal Matrix Composites", Journal of Composite Materials,
I
in press.
22.
I
S. K. Raghunathan, C. Persad, D. L. Bourell and H. L. Marcus, "High-energy, High-rate Consolidation of Tungsten and Tungsten-based Composite Powders", Materials Science and Engineering, A131 (1990) in press.
3
23.
and High Energy High Rate Processing of Ultrafine Grained Tungsten", Proceedings
3 I
H. L. Marcus, D. L. Bourell, Z. Eliezer, C. Persad, and W. F. Weldon, "Synthesis
of BTI Advanced Armor/Anti-Armor Materials Program Review, in pres 1990.
24.
3
Z. Eliezer, C. Persad, S.C. Sparks, D. Moore, M. Schmerling, J. Gully, and R. Carnes "Advanced Composite Materials for High-Performance Electrotribological Applications", paper presented at ASM International's Conference on Tribology of Composite Materials, Oak Ridge, TN, May 2, 1990 and submitted to ASM
Proceedings.
25.
C. Persad, M. Schmerling, Z. Eliezer, R. Carnes, and J. Gully, "Nanosized Structures in Cu-W-WC-C Composites for Electrotribological Applications", paper
3
submitted to Proceedings of the TMS Northeast Conference on High Performance Composites, Morristown, NJ, June 1990.
3
26.
M. J. Wang, Z. Eliezer and C. Persad, "High-Energy/High-Rate Powder Consolidation of Straight Copper-Graphite Composites for Hogh-Speed, High-
3
3
current Sliding Contact Applications", submitted to Wear, 1990.
21
4D. LIST OF SCIENTIFIC PERSONNEL
Faculty nd Staff
I3. 36. 38.
1.
H. L. Marcus
2.
Z. Eliezer
D.L. Bourell 4.
C. Persad
5.
M. Schmerling
S. Swinnea 7.
R. J.Allen
T. A.Aanstoos 9.
W. F. Weldon
1.
C.J. Lund (Grad. Res. As st.)
4.
M. J. Wang (Grad. Res. Asst.)
32. S. Raghunathan (Grad. Res. As. 33. S. J.Lee (Grad. Res. Asst.) 35. I7. A. de Jesus (Grad. Res. Asst.) 38. Y.W.Kim (Grad. Res. Asst.) 310. B. H.Lee (Grad. Res. Asst.) 3 22
H. Tello (Undergrad. Res. Asst. till May 1988, then Grad. Res. Asst.)
6.
L. Govea (Grad. Res. Asst.)
9.
G. Elkabir (Grad. Res. Asst.)
I I
3
*
I
3 1 3 1
1
11.
U. Lakshminarayan (Grad. Res. Asst.)
12.
C. Jan Hou (Grad. Res. Asst.)
13.
S. Cheng (Undergrad. Res. Asst.)
14.
V. Martinez (Undergrad. Res. Asst.)
15.
R. Rice (Undergrad. Res. Asst.)
Degees Granted
1.
G. Elkabir, PhD, August 1987
2.
Y. W. Kim, PhD, December 1987
3.
C. J. Lund, MS, December 1987
4.
S. Raghunathan, MS, August 1988
5.
S. J. Lee, PhD.,December 1988
6.
M. J. Wang, PhD., December 1988
7.
U. Lakshminarayan, MS, May 1989
8.
B. H. Lee, MS, May 1989
9.
H. Tello, MS, August 1989
10.
S. Sparks, MS, May 1990.
I 3 I I I
3
23
I I I
5.
BIBLIOGRAPHY
1. C. A. Kelto, E. E. Timm, A. J. Pyzik, Ann. Rev. Mat. Sci. 1989, 19: 527-550.
2. G. F. Taylor, Apparatus for Making Hard Metal Compositions, U. S. Patent 1,896,844,
I 1
February 1933.
3. F. V. Lenel, J. of Metals, 1955,7, 158.
4. S. Isserow, Impulse Resistance Sintering of Compounds for Armor. AMMRC-TR 73-43, 1973.
1
5. C. Persad, H. L. Marcus and W. F. Weldon, High-Energy High-Rate Pulsed Power Processing of Materials by Powder Consolidation and by Railgun Deposition, UTCMSE-87-02, March
I I
1987, Gov. Reports Index # 732,954, NTIS Code AD-A179 289/4/GAR PC A07/MF AO1.
6. J. H. Gully, T. A. Aanstoos, K. Nalty and W. A. Walls, IEEE Trans. Mag., 1986, 22, 1489.
7. H. L. Marcus, D. L. Bourell, Z. Eliezer, C. Persad and W. F. Weldon, JOM, 1987, 39, 6-10.
3
8. W. F. Weldon, "Pulsed Power Packs a Punch, IEEE Spectrum, Vol. 22 No. 3 (1985), pp 59-66.
3
9. R. Laughlin, J. Gully, M. Spann, W. F. Weldon, "Continous Duty Current Collection", Final Technical Report, Wright-Patterson Air Force Base Contract F33615-86-C2673, June, 1990.
I I *
24
I I APPENDIX A
I *
I I I I I I 3 I
I U I I I
PROCESSING OVERVIEWS
A-1.
High-Energy High-Rate Materials Processing
A-2.
Single Residency Sintering and Consolidation of Powder Metal Alloys, Intermetallics, and Composites by Pulsed Homopolar Generator Discharge
25
!IM
.
i i
M
High-energy, high-rate processing, driven by fast discharging stored energy devices, offers new potential for producing materials that are otherwise difficult to create, and for the secondary processing of materials such as those derived ,(.fromthe rapid solidification technologies.
H.L. Marcus, D.L. Bourell, Z. Ehiezer, C. Persad and W. Weldon
i
University of Texas at Austin
IThe
iinputs
Ienergy
Iof
consolidation microstructures of molyb-
FUNDAMENTALS
linked into fast, low-impenlance, pulse-
denum alloy TZM PREP powders processed in about one second at increasing energy The availability of pulsed power, up to 4,300 kJikg, high-energy kinetic energy storage Idevices offers an opportunity to perINTODCTONform a wide range of experiments, The power densitylinteraction time plot During the development of kinetic shown in Figuire I gives the regimes machinery at the Center for where the processing has the most poElectromechanics at the University of tential. Processing approaches can be Texas, it became apparent that the devices are potentially excitingsources divided between direct and indirect of energy for high-energy, high-rate techniques.
forming networks. Taken together. these pulse sources offer a variety of processing possibilities with shorttime_ pwer densitiesofup to 101V cm-over large areas, with interaction tinies of 10 - toseveral seconds. Figure 1 shows the operational fields of these pulse power systems as superimposed on a broad materials processing map. Direct Processing
iHEHR) processing. The availability of the pulsed energy sources allowed
The discussed studies were performed using a 10 MJ homopolar gen.
Direct processes rely on converting
programs welding exploratory heavy cross stions in of butt metals high-
erator the pulsed processing energyassource. The power homopolar gen-
level stored energy into and megampere electrical currents passing these
rate billet heating, powder consolidations coating of mixed materials and localized heating. The key result of the exploratory work was that HEHR processing offers several unique processing advantages, with the potential creating interesting ris notofeasily produced inother mateways.
erator is an electric machine that converts stored rotational kinetic energy into electric enerU using the Faraday effect. A low-voltage, high-cur, rent device. itisan excellent power supply for applicatis that require short-time, Coupling of high-energy he hoopolar pulses.' generatorI
currents through the materials in a prescribed manner interms of pulse time and current dtrihution, inconjunction with other proce-sessuch as application of pressure linkeTeprature.T-iedn. mation i'lT'l' concepts p:'ovide the basis for a fundamental undvrt-taiding of
Nevertheless. the technique holds its own set of processing problems that
toegailtT allows production of currnt puee widths down to tons of
the potential advantages of th direct pued iwer processing le cause of
inst be Rolved before becoming commerciaily competitive C
microsecond . Small-scale experit ments are powerd by capacitor bankth
the inherent control (i" the times to reache perature. i e ofthe order
6
26 sRNAs.
Jtsh
OF METweLsS t asstup
mber 9 7
•
of the interaction times shown in Figure 1, the duration at the temperature developed can be held to relatively short times. This regime limits the .eonversion of metastable phases to their equilibrium counterparts and can be well below that required for grain growth even in simple single-phase systems. Depending upon the method of applying the pulsed power, the temperature rise time can be from 10 microseconds to one second. By using a pulsed preheat or postheat in conjunction with the high-temperature pulse, more detailed microstructural control in the workpiece can be expected. A conceptual TTT diagram for a two-component powder system with a pulsed power processing thermal cycle is shown in Figure 2. The example represented by Figure 2 is a
single component systems with heating rate dependent phase transformations, the HEHR pulsed power processing could allow the phase transformation to proceed in a desired manner. For example, consolidation of metallic glass particulates would allow the final state of crystallization to be controlled. In a microcrystalline solid with processing-dependent metastable phases, controlled pulse power processing could permit a range of structures to be produced by controlling the processing parameters. The TTT diagram also indicates other regimes of powder consolidation potential for the HEHR pulsed power processing. The primary consideration is that the consolidation processing allows bth the introduction of the power and the application of exter-
transient liquid phase, A, leading to full densification of an A + B blend, A and B could be refractory metals with melting temperatures such that
nally applied stresses. The ability to go to very high homologous temperatures would allow deformation to take place at much lower stresses and at
properties. This secondary processing that, in many cases, will be required to obtain acceptable mechanical properties, would also be done on a cold
< TMB (e.g., Mo and W, or Nb and W). This could be considered a hightemperature brazing process where no low-temperature materials are involved. Alternatively, for a system with component B alone, it represents fast heating to just below TMil, in consolidating a powder for example. For
much higher rates. This can then lead to more complete consolidation without the that primary changes wouldmicrostructural normally take
wall basis with minimum time at temperature. Consolidation fromde-a green compact to acan castrange condition pending on the relationship between -pecific energy input and the applied
TMA
place when processing for extended times at the elevated temperature. In addition, the higher temperature could also serve as the starting point for inm
S Low Inductance Pulse Forming
Homopoar oture Generators
Cmpulsators
Networks
Nin 10 4
10108.s
102
4
to 10
100
1 P43 2
101 _A
a
DEING
8
1
1
10
E
SPECIFIC ENERGY Jcm SPPP2
SURFACE
SALLOYING S
0 1 Lu a: a.-
Idesired s
o
10
SURFACE MELTING
10
,0
G
-----. Figure 2. A TTT diagram demonstrates the potential advantage of HEHR consolidation in producing non-equilibrium structures. A two-phase powder assemblage is shown thermally cycled in a short (< 1 s)time exposure. place extrusion or hot toforging of the consolidated material enhance the
stress. This and the information suggested by Figures 1 and 2 is the basis of HEHR pulsed power processing. Figure 3 shows the fineness of strucwhich can be obtained with this method of consolidation. In Figure 3a.
prior particle boundaries are evident a compact consolidated with low energy input by a powder forging mechanism. With high energy input, a finely distributed hard boride phase becomes evident (Figure 3b). The
55-60, HEHR processing. Additionally, HEHR processing can produce densities comparable to HlP'ing and power, in terms of electrical current. be placed at the appropriate region of the processed material to obtain the properties and unicrostructure throughout the workpiece. For example, the ability to concentrate a major fraction of the enerkv into a thin segment of the material using a very short electrical current pulse. which as a very high frequency cur. rent iskin effect). allows localized phase and grain structure changes to take place without changing the bullof the material. Case hardening via skin effect experiments on 1040 steel has been demonstrated as %%ell as skin heating of copper. leading to grain growth and lower strength)
IHAROENI'G
A.S=Fo
,
'behaves
10
310
10
H
46-47, hot isostatic pressing IHIP'ing): HRC 47-50, hot extrusion; HRC
6
ANouu~,,,
-
Direct processes require that the
3:
0 10
i
.....
. [..-.-
8
D
WELDING
It,
10'DPP hot e. conversion -> reduction processing sequence. Aerogel forms are particularly well suited precursors for hydrogen reduction. Alcogels may also be reduced under slightly different conditions to produce similar particle sizes. Alcogels are an intermediate product in the critical point drying procedure for tungsten trioxide aquagels. Thermogravimetric analysis indicated the aquagel contained -15 H20 molecules of which -13 were in the gel interstices and -2 were present as water of hydration. TEM observation of alcogels at 1:10,000 vacuum-dried contiguous indicated dilution arrays particle two-dimensional consisting of monosized particles with 50 nm dimensions. that
The lower size limits of these gel particles are as yet not determined. However, if indeed these are < 10 nm, and can be converted into crystalline form without coarsening, then the extremely promising work on ductile nanocrystalline oxides may be advanced by this metal route . For example, Karch et al. have shown that nanocrystalline TiO, develops due to room-temperature ductility enhanced grain boundary creep originating in short diffusion lengths and increased interfacial diffusivity (14).
12 WO03
+ 31H?
=
3 W4 0 11+3
H2 0
V V
*V
IV
3 W4 011 + 9 11,
12 WO.
+ 9 11,0
V V
12 WO,2 + 24 H,
I
=12
W + 24 H.0
F..gl. Schematic sequence of the tungsten trioxide reduction reactions in dry hydrogen at 1123K. In the simplified sequence the
possible presence
of co-existing
non-stoichiometric cxides
ais
intermediate prodict-Q iq disreoarded.
.5In Ii-,TMpooirgaho etli ugtnpwe atce maebIyrgnrdcino ugtntixd loe.Pril diesosoI10 maeeiet I4
Consolidation Parameters
Free energy and enthalpy chang, indicate that magnesium or aluminum ma be more effective exothermal reductant. (13), however, for the metal oxide product purity demands are best met by the use of high-purity hydrogen. The standard sequence of reduction reactions is shown schematically in Figure 1.
The processing parameters used for ensccdati.n pe:.nts these initial on each of these systems were derived from parameter sets previously developed for commercial tungsten powders (16). The inherent assumption is that beyond a set electrical conductivity threshold, the pulse resistive heating under pressure would provide rapid densification. The state solid identifiable most is powder mechanism densification forging. The parameters for the gel-based w were: Applied Pressure = 60 ksi (420 MPa), Specific energy input: 5000 kJ/kg. For the W-Ni-Fe/B 4 C system they were: Applied Pressure = 45 - 60 ksi (315 - 420 MPa), Specific energy input: 3000 to 7000
The reduction procedure performed at 1123K for 3600s in dry hydrogen yields metallic tungsten. X-ray diffraction studies indicate that the crystalline oxide level is below the detection limit, However, an adsorbed oxygen layer is a potential contaminant due to the large surface area/unit mass of the reduced powders, (estimated value 3-10 m2 /g) Transmission electron microscopy confirms the -100 nm average dimension of the reduced metallic tungsten particles. Figure 2 is a TEM photomicrograph of a particles. The cluster of these individual particles appear to be single-
kJ/kg. Microstructure/Property 3 and 4 provide typical Figures overviews of the rapidly densified the in developed microstructures gel-based W and in the W-Ni-Fe/B 4 C systems respectively.
or bi-crystals. the affect which factors The post-reduction size of metallic tungsten particles are incorporated into Parsons' empirical equation (15). They include the bulk density of the oxide powder, the particle size of the oxide powder, the reduction time, the bed depth, and the reduction temperature. The refinement in microstructure through the use of 50 nm trioxide in aerogel form tungsten therefore has a strong effect on the reduction kinetics, due to the low bulk density and small particle size.
I
!
radial cross-sect ion n: Fig. 3. SEM photomicrograph of a HEHR powder convoLt:h:A. by produced specimen tungsten gel-based with an energy input of 7000 kJ/kg under 420 MPa appliud pruu:.:. ev :e.. A uniform consolidated powder particle size of -200 nm
I
47
q:-!basd The
gel-based tungsten
W(w'-.,i
-Fe) i64C
was rapidly
In
i,
the
apparent
or1iontation of tho W-Ni-Fe matrix is attr-buted to flow of the material under pressure during the transient liquid-phase-assisted consolidation. The tliroe microsropic features observed are tho B4C, W and tho W-Ni-Fe eutectic.
densified to -85 % density in a single energy pulse. This densification response may be associated with the increase in total resistance derived from the now contact interparticle numerous is The microstructure resistances. uniform and consists of particles with -200 nm dimensions. At this stage, the particle size evolution sequence from the oxide gel is: gel particle dimension = s50 nm; metallic tungsten powder particle isize= - nm;200 consolidated tungsten particle size = -200 nm.
present in the RDof composite ht high-density processed a reaction products to be the s.:: ies of complox carbides and borides, whon sufficient tnorqy is delivered to beginliquid-phase-assisted densification. The lines indexable in a 10 to 80 degree two-theta scan for the W-Ni-Fe/B 4C system after liquid-phase assisted consolidation W2 C (110), (200), (211), includes W (021), (002), (121), (102), (321), (302), FeWB (001), (112), W2 B(110), (002), (200), (211), 0 and (660),(822) Fe6W6C Fe3C (101), The Ni4B 3 (211), (410), (403), (013) . also Fe 7 W 6 (119)was 1ntermetallic observed. The elemental Fe and Ni were fully reacted. These phases were not observed in the material consolidated completely in the solid state, where thd maximum applied pressure of 420 MPa was full produce to insufficient densification.
In this simple single-phase material, the primary microstructural parameters of interest are the fineness and uniformity of the microstructure. The SEM micrograph 3 indicates that these Figure in characteristics are indeed evident in the HEHR powder-consolidation-processed gel-based tungsten. Microhardness measurements confirm this finding. The average hardness value of 390 VHNs 00 (384 - 404 range) of the -85% dense tungsten compares favorably with that of conventionally-sintered tungsten, in which values of 90-320 VHN (17). These values depend are obtained upon density, particle size, and degree of cold work. The high hardness observed may be the gel-based tungsten in attributed to the fine scale of the microstructure and its uniformity, and to the strong three-dimensional connectivity developed within the consolidated mass. Post-pulse cooling under an applied pressure may also provide a hardening contribution due to enhanced dislor,density.
Ir
SEM photomicrograph of an etched radial cross-section of a Fia.. (W-Ni-Fe)/B 4 C composite-specimen produced with an energy input of 4000 kJ/kg under 420 MPa applied pressure. Full density is obtained by the evident flow of the W-Ni-Fe eutectic matrix.
*
48
Several of the carbide-forming reactions are highly exothermic, and once triggered, these reactions which occur at multiple and distributed interface nodes can be expected to maintain the high temperature developed in the early stages of the processing. Indeed this processing cycle then takes on characteristics similar to those observed in Self-Propagating High-Temperature Synthesis of ceramic phases (18). The reaction products so derived effectively transform the
nature
of
9. D. R. Ervin, D. L. Bourell, C. Persad, and L. Rabenberg, Structure and Properties of High Energy, High Rate Consolidated Molybdenum Alloy TZM, J. Mat. Sci. Eng., A102, 1988, pp.25-30. 10. K. C. Owen, M. J. Wang, C. Persad, and Z.Eliezer, Preparation and Tribological Evaluation of Copper-Graphite Composites by High Energy/High Rate Powder Consolidation, Wear, 120 (1987), pp. 117-121.
the composite
-
material by introducing the attributes of the in-situ composites, in which the reinforcing phases are produced during processing.
11. C.Persad, S. Raghunathan, B. -H.Lee, D.L.Bourell, Z. Eliezer and H. L. Marcus, High-Energy High-Rate Processing of High-Temperature Metal-Matrix Composites, Proceedings of the MRS Symposium, Vol. 120, High Temperature High Performance Composites,Reno,NV, April, 1988.
ACKNOWLEDGMENT 12. K. C. Li and C. Y. Wang. Tungsten. 3rd.ed.. New York: Reinhold, 1955.
e thank Haskell Sheinberg of Los Alamos National Laboratories who provided the W-Ni-Fe/B 4C powders. Assistance with HEHR processing was provided by Ted Aanstoos H. Tello and and Jim Allen at CEM-UT. P. C. Trimble assisted in the structure evaluation. This research was supported by DARPA/ARO Contract DAAL03-87-K-0073 and by the Texas Advanced Research Program.
Wang. rtj. application, New York: Plenum, 1979.
13. S.
W.
H.
Yih
and
C.
T.
. metallurgy.Pren Taindgten! Souirct a
14. J. Karch, R. Birringer, and H. Gleiter, Nature,Vol. 330, (1987) p.556. Electrochem. Parsons, S. 15. D. Technol.,V.3 (9/10), , 1965, pp. 280-283.
REFERENCES 1i.Peter K. Johnson, Trends in the U.S.
16. S. Raghunathan, MS Thesis, The University of Texas at Austin, Aug. 1988.
Tungsten Industry, The Int. J. of Powder Met., Vol. 24, No. 1, 1988, pp. 73-76.
17. N. C. Kothari, Factors Affecting Tungsten-Copper and Tungsten-Silver Electrical Contact Materials, Powder Metallurgy Int., Vol. 14, No. 3, 1982, pp.139-143. 18. J. B. Holt and D. D. Kingman in for Methods Process Emergent Hiph-Technology Ceramics. eds.R. F.Davis, H. Palmour III, and R. L. Porter (Plenum Press,New York, NY, 1984) p. 167.
2. F. H. Redmore, Fundamentals of Englewood Chemistry,Prentice-Hall. Cliffs, NJ. 1979. pp. 313-314. 3. C. Persad, D. R. Peterson and R. Zowarka, Composite Solid Armatures by Pulsed Power Processing: A Novel Homopolar Generator Application in EML Technology, IEEE Trans. Magnetics, Jan. 1989, pp.404-408. 4. H. Sheinberg, Int. J. of Refractory and Hard Metals, March 1983, pp.17 - 26. 5. H. Sheinberg, Int. J. of Refractory and Hard Metals, December 1986, pp. 230 237. 6. J. Fricke, Aerogels. Springer-Verlag, 1982. 7. H.L.Marcus, D.L.Bourell, Z.Eliezer, C.Persad and W.F. Weldon, High-Energy High-Rate Materials Processing, Journal of Vol. 39, No. 12, Dec. 1987, pp. Metals, 6-10. 8. G.Elkabir, L. Rabenberg, C.Persad and H.L.Marcus , Microstructure and Mechanical Properties of a High Energy Rate P/M Processed Aluminum Alloy, Scripta Met., Vol 20,October 1986, pp. 1411-1416.
49
DEVELOPMENT AND CONTROL OF MICROSTRUCTURE INP/M TUNGSTEN BY SYNTHESIS AND REDUCTION OF TUNGSTEN TRIOXIDE GELS C. Persad, S. Raghunathan, A. Manthiram, M. Schmerling, D.L.Bourell, Z. Eliezer, and H.L. Marcus
I
Center for Material Science and Engineering The University of Texas Austin, TX 78712
Abstract Tungsten trioxide gels produced from acidic solutions of sodium tungstate have been inves~igated as precursor materials for submicron metallic tungsten particles. Temperature and pH control in the gel-producing step leads,to the nucleation and controlled growth of tungsten tioxide particles down to 50 nm size. Gels of stoichiometric and non-stoichiometric
compositions have been studied. 100 nm metallic tungsten particles are obtained through rapid reduction in dry hydrogen. These particles have been consolidated to produce dense, ultrafine-grained compacts. The consolidation is accomplished by use of high-energy highrate processing driven by a homopolar generator. The consolidation response of these laboratory-synthesized powders has been compared to that of commercial high-purity tungsten powders. X-ray diffraction and electron microscopy were employed in the structure Characterization. Microhardness and electrical conductivity were used in the property evaluation. Ultrafine grain sizes (< 2 gm) are retained in the high-energy high-rate consolidated tungsten derived from a gel base. Microhardness is higher than that of commercial wrought materials. Hardness improvement is attributed to the high degree of
I
connectivity and liall-Petch hardening.
This research was supported by DARPA'ARO Contract
DAAL03-87-K-
0073 and by the Texas Advanced Research Program Grant # 4357.
Solid State Powder P,"cssing Edited by A H C:auer and J J deBiroadillo The M:nr3. MVLAhI & Matera,Z &eCty. '99"
50
inmaduction
I
Anew P/M processing route has been employed to produce ultrafine-grained tungsten. It is obtained by integration of the chemical and the mechanical metallurgy. In this paper we focus on the developmet and control of powder particle dimensions, with particular emphasis on the preparation and characterization of the precursor species : tungstic acid gels. Tungsten trioxide is an intermediate product in the conversion of tungsten-bearing ores into tungsten powders. In its hydrated form, also called tungstic acid, it may be prepared as an aquagel. This was demonstrated by Kistler in 1932 (1). A gel structure can be converted into a solid, low-density, highly porous form described as an aerogel (2,3). Such a conversion can be achieved by a series of fluid substitutions in the solid network. An alcohol-for-water substitution converts the aquagel into an alcogel. Later, liquid carbon dioxide replaces the alcohol. When the liquid carbon dioxide is removed by critical point drying, an aerogel results. This aerogel is an uncollapsed, open-structured form of tungsten trioxide with a bulk density of less than ten percent. Tungsten trioxide in aerogel form, containing sub-particles with 50 nm dimensions, provides a highly flexible precursor for efficient reduction to ultrafine tungsten powder. Powder metallurgy approaches seeking to define fundamental material behavior require materials with well-defined and controlled characteristics. Among these are purity, uniform size, and predictable morphology. Wet chemical processing using sol-gel methods makes the preparation of such powders feasible. In a P/M processing sequence, such powders enable the production of homogeneous compacted preforms, rapid sintering at relatively low temperatures, and novel possibilities for grain growth control, when compared to powders with sizes and size distributions made by conventional methods. The solid state powder processing described in this study emphasizes the characteristics of tungstic oxide aerogel to metallic tungsten powder conversion, and the subsequent rapid consolidation of this tungsten powder by high-energy, high-rate consolidation processing.
Exnerimental Procedure Materials material for tungsten preparation in this study was crystalline sodium tungstate. The starting It was supplied by the Alfa Products Div'n of Morton Thiokol. Reagent grade hydrochlioric acid was used for acidification of sodium tungstate. Preparation of Tungsten Trioxide Tungsten trioxide was prepared in crystalline and in gel forms by the controlled acidification of sodium tungstate. Control of pH and temperature led to the synthesis of the material in gel form using 0.25N hydrochloric acid at 298K. The first step in the preparation of the tungsten trioxide gel was the dissolution of sodium tungstate in a dilute 0. 1N hydrochloric acid at room temperature. This solution was further acidified urtil streaks of the white tungstic acid gel were observed. Upon further acidification. a light yellow gel was precipitated in a sodium chloride solution. 1 he ge! was ,hen filtered and vashed in flowing distilled water. The simple chemical reaction is:
I
Na 1 WOs - 2HCI --->WO .HO
I
-
2NaC1.
number of gels were prepared using fifty-gram batches of sodium tungstate. Some preparations of the aquage! \%cre c,",nve:'ed into alccgels by an alcohol.for.water ,ubstilulion. These alcogels were later processed intco aerogels by critical point drying using carbon dioxide. Samples of the gels were .aken for particle size characterization and for :hermogravimetric analysis. The remainder was placed in alumina and in copper boats for hydro*en reducton.
51
HIdwo"kdpctip of Gels Hydrogen reduction was performed by flowing hydrogen through a50 mm diameter x 1.2 m long quartz tube containing the gel-filled boats. The tube was placed in a three-zone tube furnace operated at 1123K. A hydrogen flow rate of 40 ml/min. was maintained for Ihour to achieve complete reduction. Powder Consolidation Advances in kinetic energy storage devices have opened up anew approach to high-energy, high-rate (HEHR) powder processing of refractory metals. The processing consists of internal self-resistance heating of a customized powder by a fast, high-current, electrical discharge of a homopolar generator. The high-energy high-rate " MI in Is"pulse permits rapid heating of a conducting powder in a cold wall die. This short time at temperature approach offers the opportunity to control phase transformations and the degree of microstructural coarsening not readily possible using standard powder processing approaches. The underlying fundamental approaches to higa-energy high-rate proc.ssing have elsewhere (4). The general details of the experimental apparatus and the pulse described been characteristics have also been reported (5). The technique has been employed previously in the processing of the refractory metal alloy, Mo-TZM, and for W-Ni-Fe + B4C, Ni-Mo + MoB2, Cu-Graphite, and (Ti,Nb)3AI based composites (4-10). In this study, 50g quantities of the gel-based tungsten were loaded into insulated die cavities between tungsten electrodes. Pressures between 210 and 420 MPa were applied, and the compacts were rapidly heated by a homopolar generator pulse discharge. Disks with diameters of 25 mm were produced. Specific energy inputs ranged from 2000 kJ/kg to 7000 kJ/kg. Microstructure Evaluation
I
Standard metallographic procedures were employed in the preparation of radial cross sections of the consolidated materials. Sectioning required the use of water-jet machining. Vickers microhardness measurements were performed at room temperature to determine the degree of microstructural homogeneity. X-ray diffraction analyses were performed on aPhillips diffractometer fitted with a Cu tube and a graphite monochromator. Powder samples were held onto a glass slide with doublesided adhesive tape. Sufficient powder was used so that no signal was detected from the glass slide. A JEOL 35C SEM with EDS and WDS capability was employed for powder particle chacterization and for densification studies. A 200 kV JEOL.200CX instrument was used in the TEM studies of the xerogel and the metallic powder prodLtct. Drops of a 0.01 percent solution of the alcogel were placed on holey carbon grids and vacuum dried for gel particle size and contiguity characterization.
I
Results and Discussion Thermogravi metric analyses of aquage!s and alcogels indicate that these gels are fully converted to xerogels at 420 K. Thcre are distinct mass losses as,, ciated with evaporation of ethyl alcohol (bp = 352K). the removal of adorbed and interstit;,l water, and the hoss of water of hydration from the tunstic acid I \\ 03 . \ 1120. ( 0.33 < x < 2.0,1. Post TGA for the ,tandard pattern obtainedPeak patterns wkere compared toChemical powder X-ray diffraction Co.. Milwaukee.Wl.j. - U3- '. Aldrichwkth scattenng acid (CAS ..7783 crystalline tungstic particles. fine b\ as,sociated patterns gel broadening was evident in the
*
52
II
Figure 1- SEM Photomnicrograph of a tungstic acid aerogel. Subniicron particle clusters are evident in the highly porous microstucture. Figure 1shows an SEM photomicrograph of the tungstic acid aerosel. Individual submicron clusters are evident. These clusters are crosslinked into a three-dimensionally continuous structure. The aerogel isextremely fragile, and its low density demands careful handling to
3 3
avoid breakup.
Hydrogen Reduction The reduction procedure performed at 1123K for 3600s in dry hydrogen yields metallic tungsten. X-ray diffraction studies indicate that the crystalline oxide level is below the detection limit. However. an adsorbed oxygen layer is a potential contaminant due to the large surface area/unit mass of the reduced powders. (estimated value 3-10 m2 /8). Transmission electron microscopy confirms the -100 nm average dimension of the reduced metallic tungsten particles.
*
Figure 2.is a TEM photomicrograph of acluster of these particle%. The individual particles appear to be single- or bi-crystals. The appearance of the metallic tungsten cluster suggests that the reduction mechanism is that described by Cheney (11) as Process 1. This is a gassolid reaction process that results in tungsten particles nucleated within the precursor oxide skeleton, For these gel precursors. Process I appears to be favored over the "vaporizationreduction process". which is the competing reduction process. The detailed characterization of the reduction processes that occur as crystalline tungsten trioxide is converted to metallic tunesten is the subject of numerous studies (12-1I5). the results of which do not appear to be definitive
I I
The factors which affect the post-reduction size or metallic tungsten particles are incorporated :nw Pr~os'empirical equation I161. They include the bulk densits of the xd odr h rarticle size of the oxide po~der. the recduction time. the bed -depth. and the reduction :cenperiture. The refinement in microstruci tre through the use of 50 nm tungsten trwi-sde in ."erc"oel form therefore has a %1:0111g effe~t on (lie reduction kinetic-,. due to the lo%% hu'-k 4:t and small parnicle size. Furthentore. the ease with %hich the reducing gas flo~ks .hrough the structure. and the Ablity to .%ecp out the %ater %aor which is the gaseou% rojction product.. ippeasr to fa%or the impro~ed redjacton respxonse olthese -els.
*
53
U7
-6
4"
Iiue2 E honcorp fmtli unse atcecutrotie yhdoe reutoII13 60)o ugb cdgl Iowldto aaeev 'repoaareemiu c ()rte
nni
osiliavn xv m nt
necho hs
I,,ttseed mdItmprmtrst r~osyd wlpdf.l omri4wise Im es(7.Teihrn ,Ntpini htbodasteetcl odciiytrsod thIuerst~ etn ne rv r %Nljptd a- esfcto.Tems idniibeIh tt emiainmcin drfr,,, I5 ~
Extrapolations of experimental heating and cooling rate data indicate several fcaturc ' Of thc thermal history during consolidation. The heating rate is in the 103 to 104 K/s range, and the peak temperature is -2500 K. Tbe total time at T > 0.5 Tm is less than l0 s. This duration is largely controlled by the effectiveness of conductive heat e...raction fron the surfaces of the conselidaied disk through the electrodes to the massive (10 kg) cold copper platens.
3 3
N~lic osructure/PpEr=
y
The gel-based tungsten was rapidly denisified to -85 %density in a single energy pulse. This densification response may be associated with the increase in total resistance derived from the now numerous interparticle contact resistances. In this simple single-phase material, the primary microstructural parameters of interest are the fineness and uniformity of the microstructure. Microhardness measurements indicate property improvements. The average hardness value of 390 VHN500 (384 - 404 range) of the -85% average density tungsten compares favorably with that of conventionally-sinrtered tungsten, in which values of 90-320 VEIN are obtained (17). These values depend upon density, particle size, and degree of cold work.
I 3
The SEM micrograph in Figure 3 indicates that sufficient connectivity is developed in the I IEIR powde r-consol idation -processed gdl-based tungsten to produce plastic flow. This is evidenced by the curvature of the surface on the perirnenter of the hardness impression.The high hardness observed in the gel-based tungsten may be attributed to the fine scale of the microstructure and its uniformity, and to the strong three-dimensional connectivity developed within the consolidated mass. Post-pulse cooling under an applied pressure may also provide
3
a hardening contribution due to enhanced dislocation density.
[I0 IOATS I0Ll
Ft-ure 3- SffENI Photcomicrouraiph of a Dianiond 1yraiid Indenter impres',ion 15(02~ load in the \ -!a\r;ecro.-entIn of a consolidated tun,_,ten dA,~ The 2im:n hee2,ih-rate i~' heorefical by hi d!ka\ col..clhd.1ed toa. dinqtv (if wcid zel. :..:c fr prepared in g of;powders proc. onr.o;00.ition po'Ader *
55
3
Conclusions (1)
Tungstic acid gels containing 50 nm particles have been prepared by prcptation a tngsate.Control of pH and temperature permit-. the synthesis of a gel froodim form rather than acrystalline aggregate.
(2)
The tungstic acid aquagel produced by wet chemical processing has been converted to the aerogel form y critical point drying. This dry, open-structured form of tungsten trioxide alters the hydrogen reduction kinetics. 100 nm tungsten powder particles have been prepared by hydrogen reduction of these
(3) (4)
tungstic acid gels. The 100 nm tungsten powder particles made by this mcthod are well-suited to consolidation by pulse self-resistance heating. Tungsten disks consolidated this way have microhardness values higher than conventionally sintered material. Acknowledgment
Assistance with HEHR processing was provided by Ted Aanstoos and Jim Allen at CEMUT. H-. Tello and P. C. Trimble assisted in the structure evaluation. This research was supported by DARPA/ARO Contract DAALO3-87-K-0073 and by the Texas Advanced Research Program Grant # 4357.
References iter Chrn Expanded Aerogels," JuralQofEby.ical.Chernistly. 36
I. S.S
I2.
J.Fricke, 'Aerogels," Scientific Amnericar.. May 1988. 92-97. 3. J. Fricke, Aerogels - Highly Tenuous Solids with Fascinating Properties. J. of Non. Crysaline Soids, 100 (1988), 169- 173.
3HighI UHigh
4. H.L.Marcus. D.L.Bourell. Z.Eliezer. CPersad and W.F. Weldon. High-Energy
Rate Materials Processing. Journa! of Metals. 39 (Dc-c. 1987), 6-10.
L
5. C. Persad, H. L Marcus. and W. F. Wecidon, " High Enerp~ l;gh Rate Pulsed Power Processing of Materials by Powder Consolidation and by Railgun Dc;osmton", Frinal Technical Report 4 UT-.MSE-87.12_ .:rch 1987. Gov. Reports index 732,.954. NTIS Code AD-A 179 289/4!GAR IP A07.MF AOl. 6. D. R. Ervin, D. L. Bourell, C. Persad. and L..Raheiibere. Structure and Properties of Energy, Hfigh Rate Consolidated Sic1\bdenum A11o% *FLM. JMatSci, ._n~. A 103 (1988) 25-30.
Ild *
7. C. Persad. S. Ranhunahan. B. -H. Lee. D). L. Bourell. Z. Eliezer "Ii( -nRS, 'lg-aePocsigollg Teprt eMtl anLNarus H essingo. 20 F 1),h- merrireE Mal8. MtrixComposites,2(NR -Enrg HgWiaPrc S. ha.an .R Strife (edx). \IRS. rittsburch. PA. 1988). 23 - 28. L.Bourell. and C. Per-.ad. 8. W. Y im, Cnslidtin o Netahc Glas, Ribbon'. usmng~ Electric Drtch~rge Weldin. TransaCm~n. . Voi. 19A i9S z 1634-1638 M I -jalial
'
56
9.
K. C. Owen, M. J. Wang, C. Persad, and Z.Eliezer, "Preparation and Triboiogical Evaluation of Copper-Graphite Composites by High Energy/High Rate Powder Consolidation," -W-. 120(1967), 117-121.
10. C. Persad, B. -H. Lee, C. -J. Hou, Z. Eliezcr and H. L. Marcus, "Microstructure/ Processing Relationships in High -Energy High-Rate Consolidated Powder, Composites of Ti 3 AI + TiA", submitted to Proceedings of the MRS Symposium on High
1
Temperature Ordered Intermetallic Alloys, Boston, MA, November, 1989. !1. R. F. Cheney, " Production of Tungsten, Molybdenum, and Carbide Powders"
Metals Handbook Ninth Edition, Volume 7, Powder Metallurgy, ASM, Metals Park, Ohio (1984), 152-159.
12. K. C. Livand C. Y. Wang. Tungsten. 3rd.ed.. New York: Reinhold, 1955.
3
13. S. W. H. Yih and C. T. Wang. Tungsten: Sources. metallurgy. properties and application. New York: Plenum, 1979. 14. Tao Zheng Ji, "Production of Submicron Tungsten Powder by Hydrogen Reduction of Tungsten Trioxide," (Paper presented at the Third International Tungsten
Symposium, Madrid, Spain, May 1985), 108-112. 15. V. K. Sarin, "Morphological Changes Occuring During Reduction of W0 3 ," Journal of Materials Science, 10 (1975), 593-598.
3
16. D. S. Parsons, "The Reduction of Tungsten Oxides by Hydrogen", Electrochem. Technology, 3 (1965), 280-283.
17. S. Raghunathan, MS Thesis, The University of Texas at Austin. Aug. 1988.
3
18. N. C. Kothari, "Factors Affecting Tungsten-Copper and Tungsten-Silver Electrical Contact Materials," Powder Metallurgv Int.. 14 (Mar. 1982), 139-143.
I I I I I I I I *
57
M Iaterial, Science and Engineering. .41.1 1990; XI(rn-OEN
W
High-energy, High-rate Consolidation of Tungsten and Tungsten-based Composite Powders
5 3
S. K. RAGHLNATHANi. C PERSAD. D.L.BOCRELL and H.L.MARCUS CeteicriierIttls X-ietic-e and Engineeinag. The Vniver3irt of Tt~easn4usgin. .Ativn. TX 7.Y712 fUS.A.)
Rectit d May 14. 1990: inrceis~ed form July 2. 1991)
Abstract
into electric energy using the well known Faraday
TugtnadInse-bsdhayaly r
3
effect. It is a low-voltage, high-current 'approxi-
tugMnAae a0d al-y ardvcetaeapisih miai it-ell knoisn for Tuimately their superior mzechzanical' propmtey pu 30 foAea: vier hat te6-. aple s hi-8 henrypussfraeysht UnHowever temperatures. elevated at ~ ~~ert'ies c.I~veduinste isdificul toconoliateowig ~ High-energy. high-rate !HEHRW consolidation its melingtemeraure etylug 683Kj.The Of powders employing the homopolar generator its eiyhighmelingtempratre 3683K).The as a power source has definite advantages over .additions of smiall amounts of lots-,nelting el'cnetoa odrmtlug ehius Sincentheonl~ iowdevredalug tehrtqtime MOUnts lich, as iron. nickel. cobalt and copper. ivr e din echnqeofr thme Sinbeuth1-3 s. this fiwilitate thec powtder processing of dense heavy fesso-e thisquce to3 i pres opporuty tempoeratures. moderate alloys at the s the prrvlpe mirsrcrnasoicilsed waebe opruith sd ieqtchiiqh-current pulse beener usedcnslicite wnereti oravea recentlY for posvder consolidation. In this Paper, dated material. Secondly, because of high filn the use of a hiopolar generator as a Pow~er and contact resistances and the associated pulse wurc toconohidateselctedmngenand Joule heating, effect, the energy deposition is nungsten-based alloys is examined. Va~riouis inate cntae~tteitrpril otcs lo edt-uue inthispriclescopeats Msoa cethe rials itcre consojidated including unralloved twngahs it col di setll usnfediainca inrccerated ten, WV-Nb. W-.Vi. and tungsten heavy ('1110 W'ith ceeae thtoSSi fiuat n coldw oi boron c-arbide. The efrct of pro'ess*parwizeers rsueadsefcCW~ ~~f ~~~~~~~~~~oln otdthe compact after the current pulse is dischareed. con'ljait~nof Jitferetit alloy vsrenis isclescr.,bed InI terms of Micro~tructure atid proper,, relaAwd ae~o uc~ul consolidated usine thisloshsbe H~EHR technique. Among tum.'ups.them are NMo-TZM alloy 9". nickel-based metallic -,la-s NMetelas' '10. a Cui-graphite metal1. Introduictionl and background Matrix comr-osite I1 and Al-SiC metal-matrtx ManN researchers I-; have emled Llect nic currents as a he:at source to consolidatle rpc%%ders oi different l) svstems includine iune'cen. prmnapocde Grcen,,p.rn c'nsolidated elemental ,un,_,,ten usine current densities ranging rom _') I.. an~oiiion )et-ut, to -,-, MA1 Mn': to a maximum density of -,, 1of A ichematic iketch of the experimental se:-up the iheioretical density. Near-Eheoretical den-sities is inen in Fie. I. The backine ring was used to were achie~ed with alloy additions of 'ilianinum pre--ent the preading of the die liner, which and atanium. The procesing time %%as !11nun. A cracks durine the ap-plication of pressure and bank of d.c. 2eneraiors was the powker source for sub-sequent discharge of the pulse. Usually, this tfle.,e impulse 'intering expenments. backine ring. was made of plain carbon or stainThe Power source used in the Present r,-earch less iteel. The material- used for slee%es were aluhomolar IJ tpelhvmorolar 'ei rhe mina or borcsilicat. 21~s4 The -!ecrode, Lied n hoinopoar HP enerzzo is n tis oearh %ereusually made from ETP m.d tune :hit on rsrotational kinetic ne\ copper or '-: ztairlos stee!. Die in. eris. in zhe
I I
Upimder
*
58
form of foils or disks, were placed between the electrodeq and the compact. The diameter of the insert was equal to that of the electrodes. The inserts were molybdenum. niobium. or stainless steds foils. or molybdenum disks. The specific energy input (SEI) is the amount of inery kiojolesperkilogram of the powder mass placed between the electrodes during a consolidation process. This was calculated using oscillographically recorded current and voltage time-traces. The power. at any instant, was calculated as the product of the currcnt and the voltage. The total energy input was obtained as the integral over the total duration of the pulse discharge.
I IParticle
22 .1hateria.'s Powders from differenit sources were obtained. sizes of the powders and proportion of
each constituent in the blend are given inTable 1. Three different tunesten powders were used. Coarse powder. having a minimum particle size of 5 jum and a maximum particle size of I150 am. was obtained from the Aesar division of Johnson and Matthey Company. MA. This powder was 99.98%s pure and was prepared by a proprietary precipitation process [1 31. The particles were angular as sho%%n in Fig. 2. Niobium powder was obtained from the Aesar division of Johnson and Mlatthey Company. MIA. This ponder was manufactured by a proprietary precipitation process 1131'. The minimum particle size was 10.um and the maximum particle size was 250 *um. This powder was 99.8%1 pure and wa used in pure
niobium arnd %V-Nbelemental blends. A scanning electron micrograph is shown in Fig. 3. The second variety of powder was obtained from Los Alamos Niational Laboratories. NM.
FIXED PLATEN
BACKING RING
3
POWADER MASS
'6O0.
Ft- chermaw:
ING PLATEN
ijeram .,f thc
h:!-nn% high-raie
pv%%c.-r pr()cesNing et-tip TA~BLE I
I~6
.
Scanning
:kcccron micr)gaph of as-recei~ed ,oa3rse
tun'mn prnmder Aear .The %hiie bar represents it,"Iam,
Pirtiche %iteind caiiI,,Iion of the poders used inIhi% reirch
I..''
i'h
Iii.\ -
I
Fie,
R....
59
INU"t
I1
I5V-20 211 0.11U1S I
F cnigeeto irgaho a-eevdnoir
i.5 cnigelcrnmcorp f3-eevdfn'9.3
Fi. 3Scanin eletro micogrph o asrecevedniobum
i.scanning electron microraph ofteas-receivd n
U
iT'
powder is shown in Fig. 5. Nickel was - - ed onto this tungsten powder using the methvd of
-
__ -
A
-
reduction of nickel-nitrate-coated rung-
-hydrogen
sten 14). A precalculated quantity of hydrous nickel nitrate was mixed with tungsten powder in
.2'
-
A fourth material consisting of a pre-mix*ed
..
4 Scannim! electron mie.ograph of as-receised F-%,. ltuesn po%der LA.NL . The
I Ipurity Iconsolidated
~nie bar represents
fine
I urn,
blend of W-Ni-Fe - B.,C heavy allov with boron carbide was obtained from Los Alamos National Laboratories. NNM. Its constitution is given in
The minimum particle size %kas 0.5 urn and the maximum particle size was 17 um. The reported was 99.9%. minimum. The particle% were angular with equiaxed shape as shown in Fig. 4. This posder was deoxidized in hydrogen at t 123 K for 12 h before consolidation and was in a elose bag, \ith two ioft copper disk patches, attached to th bag ito facilitate the pa-sage of the electric pukec ,hrAouk-h the po--\der m3The third type of tungsten powder. obtained imin GTrE Product-,. was m.:tnl\ used for consolidatine- \\-Ni allo's. Thi% nosder %~as 9'49o Thc minimum particle ize %ai I umn and the: maxi-mum particle size -.%as 1- umn with most of the particles l'ine in zme range 1-1t um. A
Ipur;! *
water, thus leaving a nickel nitrate coating on the tungsten powder. The coated powder mass wa subsequently reduced in a hydrogen atmosphere a873 K for I h. The resultant alloy powdr contained tungsten coated with 1-3 wt.% metallic
60
Table 1.
2.3 C'harcerization and evaluation A JEOL S' C scanning electron microscope equipped for" ener,_,-dispersiive spectroscopy EDS and %\aselent'h-dispersise spectrto)ccop \\ DS was%used for the preliminary esaluation o~f the as-receised powdurs. microstrucrure of the consolidated compac-s and fractography. Murakami'\ etchant was used for reseAlina the microstructure. X-ray diffraction analyv-is was performed using a Phillips diffractometer nited \,\ith a copper tarzet. graphite monochromator. aind a nickel filter. The densities of the compacts %weredetcrmined uwine both an immersion technique ASME PTC '.fand lineal analsis ot
photomicrographs. The latter method was used when the compacts could not be separated tram the electrodes. Densities measured using both techniques of pycnometry were checked for accuracy and were found to be in reasonably. ged ag~reemnent with each other. Specimens were prepared for hardness evaluation according to ASTIM standard E 92-72 [151. A Wilson Tukon LR* microhardness tester with Vick-r~s indenter was used to measure hardness values. A load of 0.5 kg was applied. For fractographic analysis. square rods (2mm x 2 mm., were machined from the central sections
,
K I
.
of the consolidated compacts. These specimens were fractured by bending at room temperature. The fractured ;ufcs eeanalyzed uig scanning electron microscopy. A brief summary of the systems that are discussed with their respective properties isgiven in Table 2
U. I
4.i'~ Fg .Sann
lcrnmcorp
fadaerlcos
%.ectonof the compact A I consohdaied from deo\idized fine
tungsten pouder LANL. This compact was etched with Muorakami's reaer. The -,%hit bar represents I urn.
3. Results
.1. 1.Iicro-\tnicure 3.1.1. Pure iungsten Tungsten compacts were consolidated with either fine or coarse elemental po%%ders. A micrograph of a fine powder compact Al is shown in Fie. 6. Ab is evident from the scanning electron micrograph. some deformation occurred at interparticle contacts: however, there was no iubstantial densification. The density was measured to be 75% of th~e theoretical density. The total SEI was 41$O)kJk&'- and the applied pressure was 210) MPa. The microstructure of the as-received coarse tungsten powder isshown in Fie. 7.The gain size rangzed frem -;.umto 100 yum. The compact A2
g
Fie.
Micro%.ructure of as-recelsd coarse wnuren JrA~r h thn sd~ssNuaanX:aet The %Ahitebar represents It0 um,.
TABLE 2 Summarn of the systems studied in this rese~rch and their revpectite propertt,%
W2
At
LA\NL
_'I
4i"t
A~ R _' A
Ae'4r Adr
!141 -'I
.Iu.41 I,'N.
Act4
-'III
BjI
BIt *~
61
5
244, 'J
4014,
Wa4; processed with coarse tun!,,,,ten powder at 210 MPa and 750 Ui ka'. After consolidation, it was apparent that the ends of the compact were infiltrated by copper. The density was measured to be 98%. The compact comprised a central cast region. a particulate-transition region. and a copper-i nfilt rated end zone. rhe micrograph shown in Fig. $ depicts a magnified imagte of the cast region. which was at midradius from the center of the compact. This cast recion reveals a columnar grain structure heavy etching. The tungsten particles in the copper-infiltrated region. in Fig. 9. were sintered and surrounded by copper.
Iafter
Anoth.r compact AS, was processed with coarse tungsten powder at 210 MPa and 1800 kU ke-1. The microstructire at the center of this compact. shown in Fig. 10. revealed that tungsten was not melted and -the pores were filled with copper. The density was 921o. 3.1/2 Ti'incswni-moiibwnz illo~v Tungsten and niobium powders were consolidated usinLg several parameters. A scanning electron micrograph of the 50W-5OINb compact 13 1)processed-at 210 MPa and 10 900 kk- iis shown in Fic. 11. The lig-ht areas are tungsten particles and the darker matrix is niobium. The
lip
Ithe
Fig'. I. Nficrostructure etched %ith Slurakam's reagent
t the jempact the center of .asi revion at midradiu% trom A- roeaiing columnar and equia.xed grain -,iruciure with copper bireans. The %hiie bar represents 10 am,
-AL
'
I I
VV
'
,gv e.
ret-wls, 'T% %-.. 'm .4.1
Piz. 10 %tic--ictureot: o~mpac, AS, reealing 'intered !un~ien pir:e:!. in a oppe, matrix. The dark pha.Ne is copper and :he light phase iN:un-pten. This compac, %%as es R)urn. hirkr une'ced The '%hnc)J
Dm1
*
h..
62
v.
niobiumn appears to have flowed plastically and encapsulated the tungsten particles. The density
I
was 94% Pores were present at the W-Nb interfaces in all the compacts. Another 50%V-5ONb compact ;B21,. that was processed with copper electrodes instead of stainless steel at a lower energy level C3400 U kg-' . reached a density of 76%.
IThe
3.1.3. Tungsten-nickel alloy amount of nickel coated on GTE tungsten n ws 1 2.an 3 t.% T power tisresarc study the effect of nickel, other parameters such as pressure and discharge rate were maintained constant. Compact CI was processed with tungsten powder coated with 3wt.%' Ni, at 420) MPa and 14 250 kJ kg' The density of the compact was 98%. The microstructure of the sample revealed isolated nickel pools at the center of the compact. The microstructure of CI at a nickel-rich area is
Ishown
of the zompact, wa- found it, b-e 97% of the theeretical density. U..4.Tu'ieJwv fa+8C These compacts %%ere processed using the prealloved blend of tungiten. nickel, iron and boron carbiae. Compact Dl was processed at 420 MPa and 38 10 WJke - . The density was measured to be 97'o. The micrograph of this compact is shown in Fie. 14. This micrograph reveals tungsten particles,. and the matrix in the form of sheair-band-like striations along different directions within the compact. The EDS analysis of the matrix re'ealed a mixture of tunttsten. ',ri.-.
IA.0
"
in Fie. 12. The matrix had less nickel at
:
the edge than at the center of the compact. Another compact in this alloy system :C2 was processed with tungsten powder coated w~ith 2 w6N.The microstructure of this compact is shwnFig. 13. The denbity of the compact was
IThe
*
..
I.. . '
third compact CS was processed with
iame parameters as Cl. except the tungsten powder was coated with I wt.%, Ni. Nickel pools were absent and nickel was found to be uniformlv distributed throughout the compact. The density
Fe1
lrotutr
t:matC.rvaiglre
iner13 Mirotu we aic ionp act the pres.inTi .5 Junctched The 'hit har repreents 11)urn ..ompict
II-
:omr.ic- Ci re%%jine imnered Fi \icr.o.;ructure !Ure-tenra.% in a~ niek%:l-ncn 'ruinx The Ijrk phais~u n.und zhehghih. maawn,%. phau,., :un-c'.ien Thia' rnpact Aa- unexched Th %%hitch'Ij rcre'ert, Iiiam
*
63
F~ au. %.
.
DI rc~cain: the - 'rpact %i,;rotrtuvurc ',r". matnano - he *hca~r*"nn~~ * : tt...hed .'.th \iuraikama * Tht,%a 'fOe 1.1f1.dl
nickel and iron. The average composition of the -matrix obtained from different locations in the matrix was 85.1 wt.% W. 10.3 wt.% Ni. and 4.6 wt.% Fe. Based on X-ray spectroscopy. the phases identified were W. W:C. FeWB. W:B. Fe3C. FeW 6 C, and Ni B.. The intermetallic Fe-W,, was also observed [16]. The density was 97%. The other two compacts. D2 and D3. were processed at 210 MPa and 3740 kJ kg- 1 and
transition region. and a copper-infiltrated region with sintered tungsten particles. In the central cast region, where no prior-particle boundaries were found, the fracture was found to be totally cleavage fracture. In the particulate-transition recion. where weak bonding between the panticles was found, a crack propagation path was evident. During consolidation, there was also evidence of good plastic deformation, repacking. and reshaping of the particles, as in the case of
4000 Id kg -' respectively. These compacts were solid-state sintered since no reaction product could be identified from the X-ray diffraction pattern. Densities of D2 and D3 were 93% and 92.5% respectively,
powder forging. At the edges of the compact. the infiltrated copper matrix failed by ductile fracture. whereas the sintered tungsten particles failed by transgranular cleavage. The fractograph of the pure tungsten compact A3 processed at 210 MPa and 1800 kJ kg" '. also showed ductile failure of the copper matrix and cleavage fracture of tungsten particles. From Fig. 18. necks were found between the sintered tungsten particles. The fractograph of SOW-50Nb; B2) processed at 210 MPa and 3400 kJ kg" is shown in Fig. 19. From this figure. porosity is evident. Particles appeared to have sintered only at the asperities. Upon fracture. the interparticle contacts ruptured because of the absence of appreciable bonding. A densification map denoting the corresponding density for a particular pressure and SEI combination for different systems is given in Fig. 20.
3.2. Mechanicalpropertiesandfractograpi Hardness values of tungsten ranged from 350 to 500 H,. The Vickers hardness '0.5 kg' profile of a pure tungsten compact cross-section is shown in Fig. 15(a). The hardness profiles for the tungsten heavy alloy with boron carbide and alloy are shown in Figs. 15,b, and 151c! respectively. From Fig. 1.. it isapparent that the hardness values increase from the edge of the compact toward the center. This reveals the inhomogeneity arising from the faster cooling rates on the edges of the compacts. both longitudinally and laterally. The spread in hardness for tungsten ,Fig. 15.a, and V-NbFig. I5.c isnotasextensiveas was observed in tungsten heavy alloy "Fie. 15'bi;. Furthermore. the absolute hardness values for the tungsten compact are comparable with the tungsten components commercially manufactured '17. 18Iand the hardness values for V-Nb cornpacts match the hardness of the niobium phase. The tungsten heavy alloy was processed at significantly less specific energy. which isthought to be the reason for the large hardness gradient. The variation in the Vickers hardness numbers for coarse tungsten powder compacts. with respect to the specific energy input is shori in Fie. 16. Numbers in the figure denote the aeraee values of the hardness over the entire cross-
Compact A2 Fig. 16 was found to contain three zones. a cast region at the center of the compact. a particulate-transition zone next to the cast region. and a copper-infiltrated region with sintered tungsten particles. The sequence of events that occur during the HEHR process are deduced from the above fractograph. Upon discharge of the pulse. the temperature reached the melting temperature of tungsten 3683 K and the center of the compact. the hottest part. was cast. The next region the particulate-transition region on the fractograph %as also heated but the maximum temperature was lower than the center of
section of the compacts. From this. it can be seen that the hardness of the coarse tungten compact increases with the specific energ. input to the compact. The fractograph of coarse tungten compact A2. processed at 210 %IPa and about -4I kJ k..e is shown in Fig, . 7. There were three distinct r,,tons. a central-cat region. a particulate-
the compact. These particles were plastically forged and partly sintered. As the particulatetransition zone was being sintered. the temperalure was sufficient near the copper electrodes for 'hem to melt and infiltrate through the compact. thereby enabling the molten copper to reach the interface bet,,een the particulate-transition zone and the copper-infiltrated zone. This resulted in
IV-Nb
I
*
64
4. Discussion
I()A2
HARDNESS PROFILE
(a)
A2 HARDNESS PROFILE
200
*
100'o, .10
IRADIAL I(b)
10
5
0
.5
.
(b)
DI HARDNESS PROFILE
400
-400
z 0
a
O
>~
>
10010 .5 .6
IRADIAL (Cl
IZO
.4
.2
0
2
4
6
S -
8
(c)
BI HARDNESS PROFILE
400
300
z 300
00.-
-00,
100
100
3 A 6 2 .2 0 . .6.4 RADIAL DISTANCE FROM CENTER. mm
icx r ardln% m %Iniohium B
proffic% :or a pure wrioten A-'
65
3 -
-
AXIAL DISTANCE FROM CENTER, mm
DISTANCE FROM CENTER, mm
500,
Fi-e 14
*
DI HARDNESS PROFILE
500,
Soo
~300
.
2
AXIAL DISTANCE FROM CENTER. mm
DISTANCE FROM CENTER, mm
bungmer,
.3
a
I HARDNESS PROFILE
3 2 0 1 1 . AXIAL DISTANCE FROM CENTER. mm -. Atfl %i
i'rwn
zai bide DI~n.
uni%;n
30
*
I>
1
3
4
S
6
7
SEI iMJi'kg,
I
Fie. 16 Influence of the. pecific enerep input on a'.erae hardne%%. in VHN. for compacts conw-flidated from coarseX s tungesten powkder Aesar .
0
5
Fig. 19. Fractograph of compact B2 revealing the lack of any
~~
-
appreciable bonding bet'een tungsten particles. The arrow indicates aductile failure. The %Ahitebar represents 100 urn.
.
15000*
I
0
d ,
$ 000
*1111
lip
I
V
I.
:W
i
IN34C
?5
0 100 1SO 0 Fie. I17.Fractograph of compact A2 depicting three zones: op. central :ast: centem. particulate transition: bottom. iinited turivien particles %ith copper infiltration The 'hite harmp-wnt m.ith 100
:
Fine W
'S0 300 350 400 450
PLE RSUEII& iis PLDPESI est F~2)Dniainr~p h~ svraini applied pressure and .pecific enera) input for differen't ivstem!. Ntudied in this re earch. The density values are gisen adjaccnt to the corresponding %aluesof pressure and tpecific energy input. An astermk indicates that molten copper :nfiltratcd -he com pact extcnsisel%,
rapid
conduction of thermal energy from the
then-stntering
region. paruculate-transition leavine this region just powder forged. In the cOrppr-intiltrated re-pon. the tuncstefl partcl"s
were ,intered es ident ft om the necks povibl% enhanced b%the atctisator role played by the copper in the mai, transport of tunest en 19-21'. No distinct effect of pressure- varied from 210 NMPa to 420i M?a on the W%-.Nb compacts
was found. Howveser. the density seemed to increase with the increose in ipecific ener2% input
IFit!'
Frai:.)craph )t compact sh,-Aing duecile .iimpke rtipiure .)I pjer arro% . anti nteauev tracure ,I ur einpri~scrl.Th h~ 'rrp et 1inr)
*
66
B I and B2. For the \V-3%%t.'v 0 Ni compact :,Cl . portions of
nickel matrix were istlateca as pools. espec-ally at
the center of the compact. This may be due to the tendency of excess liquid to decrease the total number of menisci. resulting in a reduction of the surface free energy. This occurs in the early stages of liquid-phase sintering [22]. Also, the rate of nickel agglomeration is rapid above its melting temperature because liquid nickel can flow along capillaries [231. With the increase in nickel content from 1 to 2 wt.%', . the sintering mechanism changed from W-W to Ni-Ni sintering [24, 25j. Near-theoretical densities were achieved by sintering the compacts at 1773 K for I h in hydrogen atmosphere 124, 25]. When these conventional processes are compared with HEHR processing, HEHR processing achieves the same density with virtually no grain growth and the consolidation process is generally complete within 2-3 s. The tungsten heavy alloy with BC is a classic example of liquid-phase sintering and this liquid phase produces a large amount of stress resulting from surface tension 123, 26]. In addition to this inherent pressure, an external pressure is applied in the HEHR process. This hassintering". been termed The liquid-phase "'pressure-assisted phase diagrams of the quaternary C-F-Ni-W , Nand tr phase d alloys are Fe-Ni-W P-81 27; and the ternary described elsewhere. Bose et aL 291. German et al '31, and BourB. and German 3 1'report that a eutectic. guignon at temperatures slightly over 1708 K. exists in the no such W-Ni-Fe ternary system.• However, • temperature for an eutectic seems to exist at thatL"' Ni:Fe ratio of 7:3 .81. However. a eutectic does exists at about 1709 K for the Ni-Fe binarv alloy system for an Ni:Fe ratio of 7:3 321. It is important that this nickel to iron ratio of 7: 3 be maintained to eliminate any iniermetallic formation between nickel and iron. which tends to embrittle the heavy alloy by segregating along the grain boundaries, thereby resulting in intergranular fracture "t0 In the compact made %%ith \V-NM-Fe- B.C. a featureless matrix and large grains exist. as has been reported by Sheinberg .33 . The hot hardnes, of this compact seems to be 1t°,,-2)% higher than that of premium-grade con entional \\'C-Co at 10-3 K.The onset of reac'ions was at I I-1 K f'or tine po%\ders and 12"3 K for coarse powders .'-'. In compact DI. no boron carbide was found. The elemental iron and nickel powder,. ,eemed to have reac.ed fully, as %as reported preiouI.
67
[331 for identical material processed at 1473 K. From the X-ray diffraction pattern, many complex borides and carbides were found. The boron carbide seemed to have reacted completely and the resulting complex borides and carbides were homogeneously dispersed. The matrix was homogeneous with the W:Ni:Fe ratio being approximately the same at different locations in the matrix. This process has the characteristics of -self-propagating high-temperature synthesis" or "SHS'" 35]. SHS is a reaction where an exotherm is involved. thereby liberating thermal energy. The liberated heat helps sustain the reaction for a longer time without any external heat input. These reaction products were not observed in compacts D2 and D3. where the constituents seemed to have undergone solid-state consolidation. 5. Conclusions t '8 e st a ci v d uhg a H H to 98% ensity was achieved ushng a HER processing technique. There was a distinct increase in the density 2,1 hardness of pure tungsten compacts with nraei pcfceeg pt 3,There was no definite pressure effect on W-Nb i t alloys. The density seemed to increase \"as sucheavy alloy with B.C Tungsten enerevipt At' hieh consolidated. ~cessfuliv input.
.y various products leding to a -self-sustaining i l high-temperature synthesis" were identified. At this high energy level, a liquid-phase-assisted consolidation mechanism is operative. , W-Ni alloy systems were consolidated successfully using nickel-coated fine tungsten poAders. Acknimedgiens The authors wish to thank Haskell Sheinberg of Los Alamos National Laboratories for pro'idinu tungsten and preblended composite powders. The authors gratefully acknowledge the a,,istance pro\,ided %ith the HEHR process by Jim Allen and Ted Aanstoos of UT-CEM. NI. Schmerline. H. Tello. and S.Cheng assisted in the ciructural esaluation. This research was supported by DARPXARO Contract DAAL03 -8 7K-Ou-u'3.
Rercrences
II I3 I4
18 2
6 7 8 9
I
F.V. Lenel. Resistance sintering under pressure. A.Met.. .jatauary. 19531158-167. J. Greenspan. Impulse resistance iinicring of tungsten. AMM.1RC TR. 76-32. September. 1976 (Army Mfechanics, and Materials Research Centerp. A-I. Raitchenko et at..Mass transfer and homogenizing in eectrie-discharge sintering of powder mixture. In M. M. Ristic ed.. Siniering-New Devel~opments. Elsevier. Amsterdam. 1979. pp. 76-8 1. D. J. Williams and W. Johnson. Neck formation and grou~th in hich-voltage discharge forming of metal po%%dcrs. Powder .leta1t., 25 :)t1982,,85-69. 5 I.Shkey i t. Eecrca dscare owercopation. Powder.Metal!. Int., 11 .31,(19791 120-124. H. L. Marcus et at.. High-energy. high-rate materials processing. J..Ve:.. J9N 21! 1987'6- IC J. B. Walters and T. 'A. Aanstoos. Weldine and billet heating with homopolar generator. Wet. Progr.. 127 ,25-28. G. B.Grant etia.. Weld. J., 58 1979,124-26. D. R.Er~tn. D. L. Bourell. C. Persad and L. Rabenbers. Structure and properties of high-energy. high-rate
consolidated molybdenum alloy TZM. 11ter. Sit. Eng.. .4102 1988 25-30. It) Y. W. Kim. D. L. Bourell and C. Persad. High-energy. high-rate powder processing of a rapidl%-quenched quaternary alloy. Ni,.,Mo:. .Fc,,,B,,,. Vter. Sci. En.. .412:3 19)t) 99- 11. 11 M.J.Want!ecit. High-energy highi-rate consolidation of copper-graphite composite brushes for high-s~peed. highcuirrent applications. In J. H.Gully ed.. Proc. Thi:rd In:. Co ifn Curent Collectors. .4-l31n. T..Noiember. 1'*W-. Piper 20) 12 H.L. Marcus c at.. in F.C.Matthew~s. N.C.R. Buskell. J. M. Hodgkinson and J. Norton eds... Interyaices in A.ltnn"n4ni Mtat-Mlatrix Compmtiw.. Vol. 2. Proc. StA:Ii /it:. Coant. on Composite $1tertals, X0M IId ECC.V 2. Els%ster. London. 1987. pp -'.459-1.468, I3 L. B. Martin. Aesar Diision. Johnson & Matbey Seabrook. NH. pnivate communication. March I98 14 C J.Li asnd R. MI.German. Enhanced iintenne of tung,.ten-phase -quilibria effects on properties. tnt. J. Pai.der IS tamiw Po~ae of ASTI ShVna149- 1ar1
ICornpanN.
o 16 / re.all.
-
*
eTe ieci chnol.. .Y)emo
hih
%it~later Re lio teicrawurcealmti cop swilp P',,( -,11i .232 R. S'.re. H,:ndh'ok on i/hc Properties oi \iol':wn. 1lAh'niti. Tantiuli. Turgn.'en and some of Viewr
68
19 20
21
22 23 2-1 25
Alloys. NATO Advisor% Group for Aeronautical Research and Development. May. 1965. Aerospace.Sinucirat .leiats HandbooA. Department of Defense Materials Information Center. March. 1963. A. N.Pilyankevich et at.. Special feature.%of failure of the tungsten-copper sintered pseudo~iloy. Sov. Postder Metal!. tnt..,25i6iiI986 516-321. V. N. Panichkina and V. P. Smimo-. Segeeation of solutes to the grain boundaries of tungsten -copper pseudoalloys. Vot. Poi~der Metl. Int.. 15~4 1986. 331-333. N. K. Prokushev etitit.. Kinetics and densification and grain growth of refractory phase grains in the liquid phase sintering of very finely divided tungsten-copper materials. Sot-. toi~der1tatl ~tt.. 25t9t. I;.,-733 Y.S.K~on and 0. N, Yoon. The liquid-phas.e sintefing of W-Ni. Sintering Procecsses. Mterials Science Research. Vol. 13. Plenum Press-New York. 1980. pp. 2 03-2?18. R. St. German. Liquid Phase Siniefrig. Plenum Press. New York. 1985. J.S. Lee and 1.H-.Moon. The dependence of solid-state sintering of W-Ni po%%der compact on Ni concentration. Scr.Metatt.,21 1987 1175-1178. R. M. German. The effect of the binder pha.%e melting temperature on enhanced sirntering. .1letall. Trans .4. 17
1l986 903-906.
26 F.V. Lenel. Postder .1ketalturr -Principles and .4pptications. MPIF. Princeton. NJ. 1080. 27 A.Gabriel ei at.. in G.C, Kucz~nski et at. eds. .Sim.ering Ne\%York. i987. pp 379-3 3. '.V. Plenum Pre%%. S F. Guillermet and L O~ilund. Expeniment.% and theoretical study of the phisc-equilibria in the Fc-Nt-'A i.stem..Mceatt. Trans_. 1- 198e. t809-l~3 cit.Te~t temperature and strain rate effects on 29 A. Bstet the properties of tungstien heavy alloy. .- ait. Trit,\ 4. 10 1988 487-4Y.3 30) R. M. German et al, Mlicrostructure limitations of hiah tungsten content hca%\ alloys. J, Yet:. August. 19-S 36-39 31 L. L. Bourgut non and R.St. German. Sintering ternperature effects on a :ungsten heavy allo\. it,.J. Po"~aer lievVeatI. 24 2 198811-11 32 N. V. Agew ed..- Handbook of Binary ,1leiallic Systems. Vol. 11 translated from Russian. Israel Provram for Scienific Translation'. Jerusalem. 1967. 3 H. Sheinbere. Non-con'entional tungsten-base hard nslei.. r entd9r3o1--Z Ila metal bit, J. HfhireSinwnior meatn J. J4-12
r %I',-:o. Delard. .%a.1983-6
1v D Corl'n ci .d SI-IS ,li-oicrine of( rnastcrial. in -Inc titinium-hvrton-.irbn %%'tcm. In Gi Brugcec.mani .ind . 11~". .:: liter~ais 1'oteisl:'w. %ol V \\e1\,. ed% , l:z ',-46 111,Plenum Prc% , Ne\%%%6. IQtA.pp
I
U
BTI ADVANCED ARMOR/ANTI-ARMOR MATERIALS PROGRAM
SYNTHESIS
I I
AND HIGH ENERGY HIGH RATE PROCESSING OF ULTRAFINE GRAINED TUNGSTEN H.L. Marcus, D. L. Bourell, Z. Eliezer, C. Persad, Center for Mat. Sci. and Eng., and W. F. Weldon, Center for Electromechanics, Univ. of Texas, Austin, TX 78712.
3The ABSTRACT
U
High-energy, high-rate processing, driven by fast-discharging stored energy devices, offers new potential for materials science oriented developments. This is evidenced by the tungsten-based research thrust described in this work. Tungsten trioxide gels produced from acidic solutions of sodium tungstate have been investigated as precursor materials for submicron metallic tungsten particles. 100 nm metallic tungsten particles have been obtained through rapid reduction of open-structured aerogel precursors in dry hydrogen. These particles have subsequently been consolidated to produce dense, ultrafine-grained powder compacts. The consolidation step is accomplished by use of high-energy high-rate processing driven by a high current pulse from a homopolar generator. The consolidation response of these laboratory-synthesized powders has been compared to that of commercial high-purity tungsten powders. Binary systems such as W-Cu and W-Mo have also been processed by similar methods. It has been further demonstrated that Cu-W-WC nanosizd composite structures for electrotribological applications may be derived from these processing approaches. This work was supported by DARPA under ARO Contract DAAL 0387-K-0073 and by the Texas Advanced Research Program Grant # 4357.
Nov. 8, 1989, Alexandria, VA. *
69
I *
SYNTHESIS AND HIGH ENERGY HIGH RATE PROCESSING OF ULTRAFINE-GRAINED
* I I * I
~TUNGSTEN
I H. L. MARCUS, D. L. BOURELL, Z. ELIEZER, C. PERSAD, CENTER FOR MATERIALS SCIENCE AND ENGINEERING, AND W. F. WELDON,
I
CENTER FOR ELECTROMECHANICS,
*
THE UNIVERSITY OF TEXAS AT AUSTIN
I
I
SUPPORTED BY DARPA/ARO AND TEXAS ARP
I *
I I I
IDA, Nov. 8, 1989
70
I
PERFO RMANCE
*.*PROPERTIES SYNTHESIS5/ PROCESSING
I
STRUCTUREICOMPOSITION
71
U
* I I
*
OUTLINE o HEHR PROCESS
I I I I
i 1 i I I I
o W SYNTHESIS o PROCESSING > W from gel > W-Cu > W-Mo 72
I
MATERIALS
i'
I I ELEMENTAL:
TUNGSTEN: GEL-BASED (i00nm)
FINE (0.8 km)
i
COARSE
i
i
BINARY SYSTEMS: W-Cu
i
W-Ni W-Hf
W-Nb W-Mo.
I
-
I
TERNARY SYSTEMS:W-Ni-Fe
i
W-Mo-Re
I ICOMPOSITE:
W-Ni-Fe + B
I
C
HIGH ENERGY HIGH RATE
'1MJ in is'
I Conversion of Rotational KE into Electrical Energy In the form of a High Current Pulse
I Faraday Disk ---> Homopolar Generator (HPG)
I I I I I
74
I
R
CENTER
U
PERIPHERAL
I
75
Compulsators
Low Inductance
Forming 3Pulse Networks o0 1010-610-
Homopolar
Generators
10
10 0
10.-2
SPECIFIC ENERGY J/cmn
I
SCK
HARDENING 1
C4E10
10 EP
I
10
PPP2
10
2
1
060a
1
10
MLIG1
010
010~
pp
10
10I
10
10 INTERACTION TIME
*
76
10 -210
~ -SECONDS
10
LIUI
II
TM
II
TMA
I//I
"EQUILIBRIUM"
ISTRUCTURE
IU
HIG-ENRGYHIGH-RATE PROCESSING REGIME
i
< 1 SEC.
TIME
II
|
77
POWDER PROCESSABILITY
II I.
RESISTIVITY DEPENDENCE
I I FINE POWDERS
COARSE POWDERS
I I
' C"esv#. P dUUr,
-4u€
INFLENCOE OF pTAeT
I
I
I
78
S1Z. .LE
MAKING RL>> RS Ii
SIIR
I
*
HPGL 100-500 kA
'LOAD
-
I
I
s
R
ISYSTEM
7
W PROCESSING REQUIRES HIGH TEMP.
1 1 I
HIGH TEMT. ---- > LARGE I2 RL
IIS DETERMINED BY SYSTEM RL DEPENDS ON POWDER CHARACTERISTICS 79
I
I II
FINE TUNGSTEN POWDERS? 1.No commercial sources of powder < lgm (1987 -->). 2.Laboratory methods
*
-Vapor phase decomposition of WF6 & WCI6 -Freeze-drying
-Sol-gelroute
80
of APT
I
* I I. I I *
0 (b)
(a)
I I I I I (c)
I I I
(a), sol; (b), gel; tc), powder
nt/S IVY 81
olmerization behavior of silica..
I
I
MOSTWW
Monomer
4Dimer
I
Cyclic. pH< 7Particle or
pH 7-10 with
salts present
pH7-10with
//3lS I nm
5nm
bsn
salts absent
H
0
A ,nm
30 nm
100 nr
I
Three-dimensional
I
gel networks
I
|
Sol$
82
(~e?8
U
TUNGSTEN via P/M *Conve,--.ntional Method
I
ORE - TO - POWDER CONVERSION
1
Woiframite (FeMnWo4)1
~heie(a0)
I0
I IT IA
~
CRUDE__
NA2W__4_____
0
N D
ISODIUM TUNGSTAE (Pure Na2WO4)
G E
AZ D
CDHW4
ITNSI I
V
>1 jim
*
9
10nm lOPOD
83
POTENTIAL PRECIPITATION PRODUCTS
1)
TUNGSTIC OXIDE MONOHYDRATE (CRYSTALLINE), WO . H20
> 2N ACID/ 1000 C
2)
TUNGSTIC OXIDE DIHYDRATE (CRYSTALLINE), WO 3 . 2H 2 0
0 > 0.5 - 2N ACID/25 C
3) COLLOIDAL TUNGSTIC ACID, nH 0 0.33 2000 m/s
I
Velocity
> 200 m/s
*
Current Density
2 kA/cm2
1
I I
Lifetime
1000 Pulses
Single Pulse
*
Material
Cu-Gr-Sn-Pb Composite
7000 series Aluminum Alloy
I
i
I I
101
200 kA/cm2
INTRODUCTION APPLICATION OF ELECTRICAL BRUSHES
I RESISTIVE HEATING PS 1 RS 1.8 kW 2
I RESISTIVE PB
OF STRAP
HEATING OF BRUSH
12 RB =500 W
IN TERFA CE HEA TING FRICTIONAL LOSSES- Pf = ;-,Nv ELECTRICAL LOSSES
1I5 kA
V =2 VI
v 200 ms
R = 20 PD
g
RS =72 pD
0.=.2
N =45 N
3
Pe
102
:1.8
kWV
VI =10 kW
C,,
A
IL
00
I
103
W-CU ELECTRICAL CONTACTS: DENSITY & HAPDNE.S
* *
RED. CuWO4
U
Cji-COATEDW ELEMENTAL
*
INFILTRATED
I
0
60
40
20
80
Ioc
DENSITY %THEORETICAL
IREDCu4 Cu-COATEDW H U
ELEMENTAL
INFILTRATED 0
20
4'0
60
80
AVG. SUPERFICIAL HARDNESS HR1 ST (Cum40)
*
104
100
Ib Ud U4
I105
4
ELECTROTRIBOLOGICAL BEHAVIOR Cu-COATED TUNGSTEN (90W/iOCu)
T max. = 560 C. T avg. = 380 C.
I *I *
SEM: DIAGONAL WEAR TRACKS ON SURFACE AFTER HIGH-SPEED TEST.
I
(IN AIR, UNCOOLED, lOON DOWNFORCE. RUN FOR 60 s @ 100 rn/s WITH ZERO CURRENT, THEN FOR 30 s @ 100 m/s WITH 1 kA/cm2.)
*
J mm x 12.5 mm W-Cu CONTACT 20 on SMOOTH 4340 STEEL ROTOR.
I
106
I
COPPER-TUNGSTEN PLANAR COMPOSITE
10
Ca
U)
3:0 a)
CZ
c
CL
*0
C
cc~
cu
(1)~ 0
44 im) B-SiC particulates were deliberately supplied with a graphite-enriched surface layer to enhance local electrical conductivity. This layer extends 1-10 im below the surface, depending upon particle size (6). The mass of SiC was 20 percent of the total composite mass, corresponding to a volume fraction of 30 %.
*
I
W-Ni-Fe/B 4 C
The tungsten-based composite was supplied as a premixed powder blend by Los Alamos National Laboratories, NM. The metallic matrix elemental powders W:Ni:Fe were mixed in a mass ratio of 93:5:2. The mass of B4 C was 1.3 or 1.65 percent of the total composite mass, corresponding to a volume fraction of 8% or 10%. Detailed characteristics of such powder blends have been reported previously by Sheinberg [5]. Processing Advances in kinetic energy storage devices have opened up a new approach to powder processing of High Temperature Composites. The processing consists of internal heating of a customized powder blend by a fast electrical discharge of a homopolar generator. The high-energy high-rate "IMJ in 1s" pulse permits rapid heating of a conducting powder in a cold wall die. This short time at temperature approach offers the opportunity to control phase transformations and the degree of microstructural coarsening not readily possible using standard powder processing approaches. The underlying fundamental approaches to high-energy high-rate processing have been described elsewhere [7]. The general details of the experimental apparatus and the pulse characteristics have also been reported[8], and the technique has been employed in the processing of Al-SiC composites [6,9]. In this study, 50g quantities of each of the composite powder blends were loaded into insulated die cavities between copper electrodes. Pressures between 210 and 420 MPa were applied, and the compact was rapidly heated by a homopolar generator pulse discharge. Disks with diameters of 30 mm or 50 mm were produced. Specific energy inputs ranged from 2000 kJ/kg to 7000 kJ/kg. *
118
Microstructure Evaluation Standard metallographic procedures were employed in the preparation of radial cross sections of each of the consolidated composite materials. The Ti3 AI/SiC required the use of water-j t machining. Vicker'. microhardness measurements were performed at room temperature to dete -mine the degree of structural homogeneity and to aid in the detection of the re.ction zones. For the (Ti3AI + Nb)iSiC system, heat treatments of 24h to 96h at 1473K were used to follow the growth of the reaction zone. Unetched and etched cross-sections were examined in an optical microscope fitted with a Noinarski interference contrast system. This permits the polishing relief developed due to phases of different hardnesses to b come readily apparent in color micrographs[10]. From such micrographs, average zone width measurements corresponding to each heat treatment were made. These measurements are not absolute values since they are uncorrected for magnification due to random angle sectioning. X-ray diffraction analyses were performed on a Phillips diffractometer fitted with a Cu tube and a graphite monochromator. The powder samples were held onto a glass slide with double-sided adhesive tape. Sufficient powder was used so that no signal was detected from the glass slide. A TEOL 35M SEM with EDS and WDS capability was employed to verify the chemical discontinuities associated with the formation of reaction products at the matrix-reinforcement interface. RESULTS AND DISCUSSION Consolidatio, Parameters
I
The processing parameters used for these initial consolidation experiments on each of these composite systems were derived from parameter sets previously developed for the unreinforced matrix materials. The inherent assumption is that beyond a set low electrical conductivity threshold, the pulse resistive heating under pressure would provide rapid densification. The most identifiable solid state densification mechanism is powder forging. Densities greater than 95% of the calculated theoretical values were o-:ained. The parameters for the (Ti3AI + Nb)/SiC system were: Applied Pressure = 30 ksi (210 MPa), Specific energy input: 2000 to 4000 kJ/kg. For the W-Ni-Fe/B 4 C system they were: Applied Pressure = 45 - 60 ksi (315 - 420 MPa), Specific energy input: 3000 to 7000 kJ/kg.
Microstructure and Chemistry Figures 1 and 2 provide typical overviews of the rapidly densified microstructures developed in the W-Ni-Fe/B 4 C system and the (Ti3AI + Nb)/SiC system respectively. The metal matrices are well consolidated, and excellent geometric conformability to the darker reinforcing phases is evident in both systems. The reinforcement appears well dispersed.
I
(Ti3A1 +Nb)/SiC In Figure 2 from the (Ti3A1 + Nb)/SiC , which was heavily etched in HCI, preferential chemical attack occured at the SiC/matrix interface where a reaction zone was observed. Consolidation appears to have occured in the solid state with the energy input of 3200 kJ/kg under 210 M:Pa applied piessure. XRD analyses of the phases present in the processed composite reveal the likely reaction products to be TiC and TiSi 2 . The TiSi phase appears to be at higher specific meltingconsolidations. associated with the regions where localized matrix ener n energy inputs has wanoobeednthloernrgy TeTi occured not observed in the lower The TiSi 2 was occured. Phases with corresponding chemistries could be identified in the vicinity of interfaces by EDS in the SEM as shown in Figure 3.The diagonal interface separates a SiC particle (lower half of photo) from the matrix. The light-colored islands on the SiC particle surface are rich in Ti
*
119
IM
Fig.L optical photomicrograph of an etched radial cross-section of a (W-Ni-Fe)/B 4 C composite specimen produced with an energy input of 4000 kJ/kg under 420 MPa features; pressure. The macroscopic eutectic. applied are A-B C, B-W, and C-W-Ni-Fe tc ae BF
Ei:.2. SEM photomicrograph of a heavily etched radial cross-section of a (Ti 3 AI+Nb)/SiC composite specimen produced with an energy input of 3200 210 kkk , under resue id 21.. .... of the SiC microencapsulation Tight under
particles is evident.
.
• -~.
Y '.F
"
i .3. SEM photomicrograph of radial cross-section an composite (Ti AI+Nb)/SiC of aunetched :.-.!i
"
specimen inwhich localized melting has occured. The region above the diagonal interface is the matrix with dark TiC reactionI products. The lower section is a SiC particle on which islands rich Ti and Si appear as light-colored areas.
*in •:.,
M AmI00 1 80
Average apparent reaction zone width measurements corresponding to treatments of a (Ti Al Nb)/SiC composite for 24h to These 1473K. at 96h absolute not are measurements values.They are uncorrected for magnification due to random angle sectioning. _i.
.interface .heat
N 40_ o
I0520
0 0
, 20
-
40
60
80
100
HOLDING TIME AT 1473 K (h) 120
4_
and Si and have very little carbon, consistent with the XRD observation of TiSi 2). EDS analv~s of the dark phases visible above the diagonal interface show these to be TiC. The multiphase reaction zones were grown by heat treatments at 1473K for 24 to 96h.The apparent width of the reacted layer measured from a series of etched cross-sections is plotted in Figure 4. Additional unidentified phases were produced during these heat treatments. TEM investigations of the p'.:ses present ar': currently underway. (W-Ni-Fe)iB 4 C
I
In Figure 1,the apparent orientation of the W-Ni-Fe matrix is attributed to flow of the material under pressure during the transient liquid-phase-assisted consolidation. The three microscopic features observed are the B4 C, W and the W-Ni-Fe eutectic. XRD analyses of the phases present in the high-density processed composite reveal the reaction products to be a series of complex carbides and borides,*hen sufficient energy is delivered to induce liquid-phase assisted consolidation. The lines indexable in a 10 to 80 degree two-theta scan for the W-Ni-Fe/B 4 C system after liquid-phase assisted consolidation includes W (110),(200),(211), W2 C (021),(002),(121),(102),(321),(302), FeWB (001),(112), Fe 3 C (101), Fe 6 W 6 C (660),(822) and W 2 B(l10),(G02),(200),(211), Ni 4 B3 (21 1),(4i0),(403),(013). The
intermetallic Fe7 W 6 (119)was also observed. The
I
elemental Fe and Ni were fully reacted. These phases were not observed in the material consolidated completely in the solid state, where the maximum applied pressure of 420 MPa was insufficient to produce full densification. Several of the carbide forming reactions are highly exothermic, and once triggered. these reactions occurring at multiple and distributed interface nodes can be expected to maintain the high temperature developed in the ear!*" stages of the processing. Indeed this processing cycle then takes on characteristics similar to those observed in Self-Propagating High-Temperature Synthesis of ceramic phases [11]. The reaction products so derived effectively transform the nature of the composite material introducing the attributes of the in-situ composites, in which the reinforcing phases are produced during processing. As a clearer understanding develops of the complex physical metallurgy of the high-temperature matrix materials, in particular the new RSP aluminides, alternative approaches such approaches crucial. guidelines selection of reinforcement to the been effort in beingSome applied to our for continued by Fine etphases al [12],become and are have discussed high-temperature metal matrix composites.
I
ACKNOWLEDGMENT We thank Ralph Anderson and Sand) Shuleshko of Pratt and Whitney, Read Stewart of Superior Graphite, and Haskell Sheinberg of Los Alamos National Laboratories who provided
I
the powders. Assistance with processing was provided by Ted Aanstoos and Jim Allen at CEM-UT. Mike Schmerling, H. Tello, H.-G. Chun, C.J. Lund, and Ming-Jy Wang assisted in the structure evaluation. This research was supported by DARPA/ARO Contract
DAAL03-87-K-0073.
121
I
REFERENCES 1.
P. K. Brindley in Hi-h-Temperature Ordered IntermetjI~c \P .e iited by N. S. Stoloff, C. C. Koch, C. T. Liu, and 0. Izumi (Mate.- Res. Soc. Proc. Th1Pittsburgh, PA 1987) pp. 419 -424. 2.
I
I I I
3.
A. G. Metcalfe in lnte-rfaces in Metal Ma trixCm*j~ Volue 1 (ed.),( Academnic Press, New York, NY, 1974) pp. 67-123.
.C
ecl
C. G. Rhodes, A. K. Chosh, and R. A. Spurling, Metallurgical Transactions A,
Volume 18A, Dec. 1987, pp. 2151-2126. 4.
I-I. Sheinberg, Int. J. of Refractory and Hard Metals, March 1983, pp. 17
5.
H. Sheinberg, Int. J. of Refractory and Hard Metals, Decemzer 1986, pp. 230 - 237.
6.
G. Elkabir, PhD Dissertation, The University of Texas at Austin, August 1987.
7.
H. L. Marcus, D. L. Bourell, Z. Eliezer, C. Persad and W. F. Weldon, Journal of Metals, December 1987, pp. 6- 10.
8.
G. Elkabir, L. K. Rabenbergv, 20 (10),1986 pp. 1411-1416.
9.
H. L. Marcus, L. K. Rabenberg, L. D. Brown, G. ElkabIr, and Y. NI. Cheong, in ICCM VI/ECCM 2-Composit Materials, eds. F. L. Matthews, N. C. R. Buskell. J. M. Hodgkinson and J. M,,orton ( Elsevier Apolied Science, London, 1987) p. 2.459.
10.
C.E. Pnce in Metals Handbook. 9th edition., V .S.-.Metao ra hv and Mi c r )tr[u ctiLr
-
26.
C. Persad and H. L. Marcus, Scripta Metallurgic a.
(ASM, Metals Park. OH, 1985) p. 15 1. 11.
J. B. Holt and D. D. Kin-man in Emer.gent Process Methods fo Hioh-Techn I-~ Ceramics, eds.R. F.Davis, H. Palmour 111, and R. L. Porter (Plenum Press, New
York, NY, 1984) p. 167. 12.
M. E. Fine, D. L. Bourell, Z. Eliezer, C. Persad, and H. L. Marcus , submi-itted to
Scripta Met., March 1988.
122
MICROSTRUCTURE/PROCESSING RELATIONSHIPS IN HIGH-ENERGY HIGH-RATE CONSOLIDATED POWDER COMPOSITES OF Nb-STABILIZED Ti 3 Al + TiA! C. Persad, B.-H. Lee, C.-J. Hou, Z. Eliezer and H.L. Marcus, Center for Materials Science and Engineering, The University of Texas at Austin, Austin, Texas 78712. ABSTRACT A new approach to powder processing is employed in forming titanium aluminide composites. The processing consists of internal heating of a customized powder blend by a fast electrical discharge of a homopolar generator. The high-energy high-rate "IMJ in Is" pulse permits rapid heating of an electrically conducting powder mixture in a cold wall die. This short time at temperature approach offers the opportunity to control phase transformations and the degree of microstructural coarsening not readily possible with standard powder processing approaches. This paper describes the consolidation results of titanium aluminide-based powder composite materials. The focus of this study was the definition of microstructure/processing relationships for each of the composite constituents, first as monoliths and then in composite forms. Non-equilibrium phases present in rapidly solidified TiAI powders are transformed to metastable intermediates en route to the equilibrium gamma phase. The initial single phase beta in Nb-stabilized Ti 3 AI is transformed to alpha two with an intermediate beta two phase. In composite blends of TiA! Powders mixed with Nb-stabilized Ti 3AI powders a 10 pm thick interfacial layer is formed on the dispersed TiAI. Limited control of post-pulse heat extraction prevents full retention of the rapidly solidified powder microstructures. INTRODUCTION The Ti-Al system offers a rich variety of phases and microstructures. Many of its equilibrium phases are disordered solutions with wide composition ranges, or simple ordered structures based on the disordered solutions (1]. Non-equilibrium processing such as the production of powders and ribbons by rapid solidification alters the phase boundaries 12,31. Alloying additions further complicate the precise prediction of phase fields and it becomes practical to adopt the use of a qualitative Ti-Al pseudo-binary [4]. At the focus of alloy design and development effort remain the binary intermetallics: TEAI (gamma) and the Ti3 AI (alpha two). Lipsitt has reviewed the history, progress, and potential uses of these materials in aircraft turbine engines [5,61. The TiAI intermetallic is of lower density (3.76 vs 4.15- 4.70 gm/cm 3 ) but is also less ductile at room temperature (1-2 vs 2-5 percent elongation) than the Ti 3AI (5) . The density advantage of TiAI is enlarged when the influence of common ternary additions such as Nb, Mo, and Ta to Ti 3AI is taken into account. Powder-based composites of Ti3 AI + TiAi have therefore been designed to take advantage of the ductility of the Ti3AI as a matrix and the low density of the TiAI as a dispersed constituent. In addition to the increased specific strength and potential thermal stability, these composites attempt to exploit the lattice parameter matching principles espoused by Fine et al. 171. The phases of interest for these selected constituents are alpha, alpha two, beta, beta two, and gamma in the pseudo-binary Ti- Al phase diagram. The beta two phase is stabilized by the presence of Nb [1,4,81. The characteri,tics of the phases are indicated below. alpha alpha two beta beta two gamma
A3 DO19 A2 1D03 LI 0
Mal Me Soc SymD Proc Vol
I
hexagonal-close-packed (hcp) structure ordered phase, based on hcp structure body-centered-cubic (bcc) structure ordered phase. based on bcc structure ordered phase. based o! fcc structure
133 *IM1 Materal,
esearch Socmolv
123
II EXPERIMENTAL PROCEDURE
i
M au:Rapidly solidifie itanium aluminide powders (-80 mesh) were supplied by United Technologies Pratt and Whitney Division , FL. As-received TiAI powders (51 aL% Ti, 49 at.% A) and Nb-stabilized Ti3 AI powders (65 at% Ti, I I at% Nb, 24 at% Al) were used. Pfocming: The processing consists of internal heating of a customized powder blend by a fast electrical discharge of a homopolar generator. The high-energy high-rate "IMJ in Is" pulse permits rapid heating of a conducting powder in a cold wall die. The underlying fundamental approaches to high-energy high-rate processing have been described [9]. The general details of the experimental apparatus and the pulse characteristics have also been reported [10,11]. In this study, 50g quantities of powders each of three groups of materials were processed. Group I consisted of TiAl powders. Group II consisted of Nb-stabilized Ti3AI powders. Group III consisted of composite blends containing 10, 20, 30, 40, or 50 mass percent TiA! powders mixed with Nb-stabilized Ti 3AI powders (+ 200 mesh). The powders were loaded into insulated die cavities between copper electrodes faced with molybdenum foils. A pressure of 210 MPa was applied, and the compact was rapidly heated by a homopolar generator pulse discharge. Disks of 50 mm diameter were produced. Specific energy inputs ranged from 3200 kJ/kg to 4000 kJ/kg. Microstructure Evaluation: Standard metallographic procedures were employed in the preparation of radial cross sections of each of the consolidated materials. Disk sectioning by water-jet machining minimized mechanical damage. Vickers microhardness measurements were performed at room temperature on both the as-received and consolidated materials to indicate the influence of structural change# on mechanical properties. For the consolidated T3AI,heat treatments of 24h at 1173K were used to determine the equilibrium microstructure. Unetched and etched cross-sections were examined in an optical microscope fitted with a Nomarski interference contrast system. This permits the polishing relief developed due to phases of different hardnesses to become readily apparent in color
micrographs 112). X-ray diffraction analyses were performed on a Phillips diffractometer fitted with a Cu tube and a graphite monochromator. The powder samples were held onto a glass slide with double-sided adhesive tape. Sufficient powder was used so that no signal was detected from the glass slide. A 200 kV JEOL-200CX instrument was used in the TEM studies conducted on fully dense disks sectioned and thinned by electropolishing [13] for observation. RESULTS AND DISCUSSION Consolidation Processing & Microhardness: The processing parameters used for these consolidation experiments were derived from previously developed densification maps for each powder. Densities greater than 95% of the theoretical values were obtained. The most identifiable solid state densification mechanism is powder forging. For the composition-optimized Ti-Al-Nb alloys, the best combination of strength and ductility by ingot metallurgy isobtained by forging above the beta transus and controlled cooling to produce a fine Widmannstatten microstructure (14]. Controlled cooling is also crucial to the production of crack-free Nb-stabilized Ti 3AI disks by high.energy hgh-rate consolidation. Like many other single-residence powder consolidation processes which attempt to maintain RS microstructures, the necessary large-strain deformation processing for m.etallurgical bonding and limited control of post-densification heat extraction can hirder both the desirable near-net-shape processing (15], and the maintenance of isotropic, homogeneous microstructures. Table 1 compares the average room-temperature hardnesses for powders and consolidated materials. Inevery case the processed material shows a hardness increase of about ten percent. Material-specific microstructural change,, which occur in response to the applied thermomechanical processing cycle account for these cl anges. These are discussed below.
124
Table 1: Comparison of the room-temperature microhardnesses for powders and consolidates. Comp.Inlerf aceI I
Ti3AI in Comp.
TLAJ inComp.
I
TOMA Compact TGAI Powder Thi~ CompactI TiAI Powder 0
200
400
600
Average VHN
Microstructure and Chemistry Group I : TiAI powders and consolidated monoliths: The as-received TiAI powder particles are largely two-phase. X-ray diffraction indicated that tht y were composed of the ordered TiAI (gamma) phase and the ordered Ti 3AI ( alpha two) phas :.This result is in agreement with those of Huang et. al. [3] for rapidly solidified ribbons of sitailar composition. A third contaminant phase, TiC, was also present. The relative ratios of r13Al: TiAI, judged from X-ray peak intensities, varied from 5:4 in the coarse +200 mesh (>74 gm) powders to S:lin the -325 mesh ( 2000 m/s and current densities of 200 kA/cm 2 are desired. For these applications, self-lubricating composites with high electrical conductivity are being developed. Binderless copper-graphite powder composites are promising materials for low-wear sliding electrical contacts at velocities of -200 mi/s and at current densities > 1 kA/cm 2. pulse resistive heating These materials are made by a process which employs "1MJ in Is" of a powder mass under pressure. Composites processed in this way have higher densities, hardnesses, and electrical conductivity than conventional commercial materials containing low temperature binders. The wear rate of the copper/graphite-on-steel couple at high sliding speeds (> 100 m/s) appears to be controlled by the size, volume fraction, and orientation of the graphite. For solid armatures in electromagnetic accelerators, the state-of-the-art system consists of a high conductivity aluminum alloy sliding on copper. At velocities of 1000 m/s to 2000 m/s, this type of armature undergoes a hybridization process in which a fluid is formed at the sliding interface. Energy dissipation at the interface is increased as a consequence, and the net result is a decrease in efficiency of the conversion of electrical energy into kinetic energy. For this application, a variety of composite materials approaches are under development. These include an aluminum-alumina composite with graded electrical resistance; a copper-tungsten planar composite; a tungsten-copper powder composite; and a tungsten-tungsten carbide-carbon-copper composite with nanosized constituents.
173
Introduction
In this paper we describe several advanced composite materials being developed for current collection in two classes of high-performance electrical machines: homopolar pulse generators; and electromagnetic launchers. Current collection is performed under sliding conditions hence the use of the descriptor "electrotribological.". Homopolar pulse generators transform rotational kinetic energy into electrical energy utilizing the Faraday effect. A low voltage, high impedance, megampere current source is obtained by rotating a metal disk in a magnetic field and harvesting the energy pulse during the braking of the disk to a stop. Pulsed power technology has enjoyed a recent resurgence and has been identified as a critical technology for the USA (1). The homopolar pulse generator technology is one of the enabling technologies for efficient pulsed power. More directly pertinent to this study is the recognition of homopolar generator current-collecting contacts (lnheab)
as a tribo-materials requirement for critical applications (2).
Electromagnetic launchers are capable of delivering high kinetic energy payloads such as those required for earth-to-space launch. Velocities >1 km/s are obtained by accelerating aisoidarmatur between two conductors (or rails). Barber (3) has defined this as the hypervelocity sliding contact regime and has described some of the armature/rail interactions leading to the formation of a fluid interlayer. The fluid interlayer initiation has been reported to be a strong function of both the sliding velocity and the cnacing materials (4). The current density and velocity associated with the homopolar pulse generator and electromagnetic launcher applications are mapped in Figure 1.Their schematic operational arrangements are shown in Figure 2. It is well established that a number of composite materials systems exhibit the characteristics necessary for high performance sliding electncal contacts. When based on
174
powder constituents, suclt-composites provide some of the necessary product properties by selection and assembly of material at the microscopic level. The relationship between the tribological behavior of such contacts and their basic materials science is the subject of a large body of studies (5-15). A key issue is the value of the maximum temperature that can be sustained by the contact at the sliding interface, while maintaining the efficient energy transfer associated with solid-to-solid contact. In these applications, the complex electrotribological environments encountered provide multiple interface heating sources, as indicated in figure 2 (a) for a homopolar generator brush. Achieving the necessary performance demands novel approaches to the design, processing, characterization and laboratory testing of suitable materials.
Approach A qualitative summary of the approach that we describe here includes the following steps:
- Design a material that can sustain maximum temperature at the sliding interface;
- Incorporate constituents that can provide the necessary physical properties: electrical conductivity, thermal conductivity, hardness, lubricity,etc.;
- Satisfy the technological demands of processability in the appropriate size and shape and at minimum cost.
The design of the new contact materials begins with the question: how do these materials normally fail? The simple answer is that they fail because the temperature at the sliding interface is too high. The high temperature is caused by localized dissipation of electrical and mechanical energy at the interface. The underlying theoretical limits ha,.'e been described by Kuhlmann-Wilsdorf (16). How best to handle the high interface 175
temperature? Three possibilities are: eliminate the source; actively cool the contact: design a material that can survive the elevated temperature. The thrust of this work is to design materials that can survive the elevated temperatures developed in electrotribological duty.
Experimental Procedure
I.
Materials Systems
A summary of the materials systems discussed in this paper is given in Table 1. Table 1. Materials systems being investigated for high-prformance electmtribologial applications. Contact Type! Constituents Composite
Development Status
Performance/Charact-erist ics
Commercial Faterial
Used in present generation of pulse homopolar machines. Fails at 220 m/s due to
Brushes
Cu-Gr-Sn-Pb
catastrophic wear. Proposed for use in future pulse homopolar machines. Hardness and conductivity exceed that of commercial material. ----------------------------------------------------------------
Binderless Cu-Gr
Composite W- Cu(90/10)
W-V;C-C-Cu
Experimental Material
Armatures Experimental
Hardness and conductivity
Material
exceed that of the 90W1OCu commercially available material.
Experimental Material
Nanosized consti--en-s increase population of potential contact spots.
I ., the solid armatures, the fabrication of a powder-based aluminum-alumina composite and alaminar copper-tungsten composite were described previously (17). The larrlinar 176
copper-tungsten composite was constructed from copper foils and tungsten mesh. These materials were stacked such that the controlled variation in spacing between the exposed tungsten fibers in the consolidated composite would yield a graded electrical resistivity.
II.
The Consolidation Process Two of the experimental materials are listed in Table 1 (the binderless copper-graphite
for composite brushes and the tungsten-copper for solid armatures) were consolidated from powders using a single-residence materials processing technique developed in our laboratories (18). This high-energy high-rate processing technique has been successfully applied to a wide variety of single phase and composite materials (19-25). A 10 MJ homopolar generator (HPG) dedicated to industrial applications was the pulsed power source used in conjunction with a modified 100 ton vertical axis hydraulic press. The dies and starting materials were initially at ambient temperature. Consolidation was performed in laboratory air at pressures of 300 to 400 MPa. After compacting the precursors to a desired pressure, a large current pulse was used to heat and consolidate the composite material. Following the discharge, the pressure was maintained for a brief holding period, thus permitting conductive heat transfer through the plungers to massive copper platens. The die was then unloaded and the consolidated compact was ejected. The voltage drop across the compact and the magnitude of the current were recorded during the pulse. The product of these values was used to obtain the power curve during the pulse. The power curve was numerically integrated to produce the total energy input to the specimen. Division of the total energy input by the mass consolidated yields the specific energy input (SEI) in units of J/g. The SEI approximates the enthalpic heat input and provides data for finding an upper bound estimate of the bulk temperature reached during processing.
177
III. Evaluation
A. Properties and Microstructure
Density, hardness, and electrical resistivity measurements were used in the preliminary evaluation of these composites. Selected consolidated composite materials were sectioned in the radial direction, metallographically prepared, and then examined by scanning electron microscopy. Image analysis techniques by gray scale differentiation of the electron image were used to quantify the area fraction of graphite in the copper graphite brushes. B. Tribological Evaluation Low Spsed Testing. The low speed tests were conducted in air on a pin-on-disk machine. Friction force and pin wear were monitored by a strain ring and a linear-displacement transducer respectively. An Apple II computer stored and plotted the data. The pin wear rate and friction coefficient were calculated from the plots. The applied load range was 2 to 20N, and the sliding velocity range was 220 to 670 mm/s. High S~d Testing. Figure 3 shows the sliding electrical contact tester (SECT). This is a device which allows rapid evaluation of materials and cooling methods with respect to interface power dissipation and wear rate. Continuous testing at speeds of up to 180 n/s at 6000 A is possible. Electromagnetic displacement probes monitor brush wear, and a video thermal imaging system provides a time-temperature history of the slider, sliding interface, and rotor. The tests were performed dry in air. No active cooling was used. Constant down force of ON was maintained during the temperature excursion tests. The specimen surfaces were prepared by polishing with a final polish using 600 grit silicon carbide paper. Prior to the test the contact and rotor surfaces were cleaned with high-purity ethyl alcohol.
178
Results and Discussion
I. Composite Brushes A.
Processing
The commercial composite is made by a proprietary method using conventional powder metallurgy techniques. No details of the processing are available. The binderless copper-graphite is produced by the high-energy high-rate technique (23-25). The degree of homogeneity of the distribution of the graphite in the copper matrix depends upon the mechanical mixing that is achieved prior to loading of the powder mixture into the die. The degree of mixing depends upon the relative densities and flow characteristics of the loose starting powders. Since the single-residence processing occurs in the solid-state, the pressure and temperature conditions within the die during the processing control the degree of densification. The copper matrix is powder-forged to high de,,,ities, while the graphite undergoes intraparticle shear to accomodate the plastic flow of the matrix. Heating rates of 100 K/s to 500 K/s are used to rapidly bring the copper matrix to the warm processing range (0.4 to 0.6 Tm, where Tm is the melting temperature of copper = 1356K). While at temperature the composite densities under an applied stress that is in excess of the flow stress within the powder volume.
B.
Microstructure and Properties
The microstructures developed in these two material systems reveal the differences in their processng. Figu. 4 shows the Cu-Gr-Sn-Pb material. The composition (in mass percent) of COM-A is 82 Cu, 11 graphite, 5 Sn, and 2 Pb. The tin was found to alloy 179
with the copper to form bronze. The lead formed small dispersoids which probably act as a low-temperature solid lubricant. There is virtually no interfacial strength between the matrix and graphite, resulting in very low hardness and low toughness. The bronze matrix helps increase the hardness somewhat, but it also decreases the conductivity. Figure 5 shows the microstructure of the binderless copper-graphite composite consolidated by high-energy high-rate processing. Low magnification shows that the graphite is fairly well distributed, but a strong orientation has developed. Higher magnification shows the tight microencapsulation of the graphite in the copper matrix. The measured values of density, hardness and electrical resistivity of the composite materials for brushes are listed in Table 2. Table 2. Properties of composite materials for brushes.
Density (gcm3) (percent theoretical)
Cu-Gr-Sn-Pb
Binderless Cu-Gr
5.5
6.30-6.68
82.2
94.1-99.8
45
62 (Axial)
Average Hardness * (Rockwell 1ST)
50 (Radial)
Electrical Resistivity
37
(g ohm-cm at room temp.) *
3.82 (Axial) 8.95 (Radial)
On this scale a fully dense copper powder compact has an average hardness of 70.7.
180
I
I II.
Composite Armatures A. Processing Cop= r-ungsten. Copper-tungsten composites exploit the high temperature hardness
and refractory nature of tungsten in combination with the ductility and electrical --
conductivity of copper and have been studied for hypervelocity contact applications such as the solid armature described earlier. In this research, four P/M processing approaches for these composites have been pursued: Cu infiltration of porous W; W + Cu blended elemental composites; Cu-W from H-reduccd copper tungstate; and Cu-coated W
I
powders, each processed by high energy high rate consolidation.
W-WC-C Nanosized Structures. Nanosized structures based on tungsten/tungsten
I
3
carbide/graphite/copper systems are being developed to meet high-performance electrical sliding contact needs. These structures provide the needed abundance of currentcollecting, refractory contact spots at the sliding interface. Sol-gel synthesis, large-strain
deformation processing and high-energy high-rate powder consolidation are the main steps in the preparation of these novel microstructures.
I
The electrotribological potential is optimized by embedding a hard, wear-resistant tungsten or tungsten carbide layer into an electrically and thermally conductive copper substrate, while retaining a graphite surface layer to act as a boundary lubricant. Tungsten oxide gel precursors allow the development of a nanosized grain structure (-100 nm) in metallic tungsten (21). This creates a large amount of a-spots to act as conduction points on the sliding contact interface. The production of the tungsten oxide aerogel is preceded by the production of a hydrated tungstic oxide gel. This occurs by the
I I
following sequence: 1.
Dissolution of sodium tungstate in dilute HCI: Na2WO4 + 2HC = W03 • H20 + 2NaCl;
I
2.
Hydrated tungstic oxide precipitates as an aquagel; 181
3.
Alcohol substituted for water to form alcogel;
4.
Alcohol removed using critical point drying to form aerogel.
This tungsten trioxide aerogel has the potential of undergoing reduction and carbothermic reactions to form metallic tungsten and tungsten carbide. A thin layer of such a material embedded on a conductive (copper) substrate would form an ideal sliding contact. Therefore the tungsten trioxide gel is mixed with colloidal graphite to form a 4:1 carbon-tungsten ratio. This mixture is coated onto the inside of a clean copper tube. The tube is next rolled flat to form a sandwich stru:ture with the carbon-tungstic oxide mixture coated inside the flattened copper tube. The longitudinal tube edges are slit and the sandwich is hydrogen annealed at 8001C for 1 hour to convert the tungstic oxide to tungsten. The copper/tungsten sandwich is then placed between two electrodes connected to a HPG. Pressure is applied and a peak current flux of 50 kA/cm2 is passed through the sandwich. This embeds the tungsten/graphite layer into the copper.
B. Microstructure and Properties CQnr-Tungsten. Results indicate that the density, hardness and electrical conductivity
of the Cu-coated W consolidates are higher than those produced by the other three approaches. The density and hardness values are higher than those for commercial material produced by conventional sintering (1700K/ I hour in hydrogen). The values are given in
Table 3. A typical microstructure developed in the consolidated material is shown in Figure 6. The lighter phase is tungsten which has a wide particle size distribution allowing for efficient filling of space. The copper, though the minority phase, is the continuous phase and is responsible for the high level of plastic-flow induced densification.
182
I
Table 3.Properties of 90:10 W:Cu composite material This Process
Commercial Material
Density
14.8 g/cm3
14.0 glcm3
Hardness (HRB)
81
70
Electrical Resistivity
11.2 p&-cm
n.a.
I I
I
III. Electrotribological Performance
An effort to optimize contact performance requires an understanding of the pertinent materials science concepts, particularly as they apply to the control of processing. The powder metallurgy based composite materials that are described in this study have been prepared by novel processing routes. Electrotribological behavior of three different systems are described: (1) Copper and Graphite (2) Copper and Tungsten (3) W/WC/C and Cu
I
I I
layered composites.
Composite Brushes
In the past, unidirectional graphite fiberstcopper matrix composites have been extensively studied for use as electrical brushes. Although their behavior in the absence of current proved to be satisfactory, t..e passage of high currents increased the friction coefficient and the wear rate to unacceptable values; apparently, the lubricating properties of the graphite fibers are inferior to those of bulk graphite. One possible reason for that may be the desorption of water vapor from the graphite fibers at high interface temperatures (26).
I I
183
Preliminary results of the tribological characterization of the binderiess copper and graphite powder-based composites were previously published by Owen et al (23), Wang et al (24), and by Eliezer et al (25). The low speed pin on disk tests (0.6 m/s) of the copper and graphite on 4340 steel showed friction coefficients in the range 0.133 to 0.207 versus
an average value of 0.120 for the commercial Cu-Gr-Sn-Pb composite. At high speeds the wear of the binderless composite is significantly lower than the commercial material. At 200 m/s values of 0.2-0.3 x 103 mm 3 /Mm versus 2 x 103 mm 3/Mm, are obtained over 6 km sliding distances. Everett (27) reported on the pulse duty ( Slkm/s), high current density (> 105 A/cm 2 ) sliding interface are characteristic of the operation of these armatures. In railguns, the residue left by the traveling armature influnces the multi-shot performance possibilities, making an understanding of wear phenomena crucial. Wear phenomena are described for three multi-leaf graded resistance armatures. These armatures were comprised of Cu-Mo-Ti (6-4), Cu-Mo-Ti (CP), and Cu-Al-Ti (CP). Each of these armatures were sliding against flame sprayed molybdenum.(Mo) on copper (Cu) rails.
3
For a sucessful railgun shot, the bore dimensions must be maintained within close tolerences. Through an understanding of the wear mechanisms occuring during a railgun shot, the wear rate of the sliding couple can be reduced. A reduction of the wear products could make the multi-shot armature railgun a reality. At present, wear debris makes reconditioning of the bore a necessity after only a few shots.
2
I I I I
233
I U INTRODUCTION
Before railgun technology develops its full potential, several problems must be confronted. One of the most perplexing is the problem associated with the rail-armature sliding couple. The armature may be considered as a single-duty high-speed sliding
3
electrical contact. However, even as a single-duty unit, the wear and subsequent wear products associated with the sliding couple must be minimized. Wear products alter the
bore dimensions and render insulators ineffective. This reduces the operating efficiency of the gun. The wear process necessitates the refinishing of the bore to remove metallic wear products from the insulators, restore uniform bore dimensions, and custom fitting of an armature to the new dimensions. The design and fabrication of these armatures to reduce wear and increase performance are the subject of an ongoing research program at The
University of Texas at Austin.
I
SOLI
ARMATURES
The primary goal of an armature in a railgun is to propel a projectile. The solid armature propels the projectile through direct mechanical stresses. The armature is subjected to high current densities ( >105 A/cm 2 ), high mechanical stress and severe resistive heating. The solid armature, as opposed to a non-solid armlature, possesses many desirable attributes. The armature-projectile pair can be integrated into one unit. Solid armatures have been predicted to be more efficient at velocities less than 10 km/s. They also have shown lower rail wear than non-solid armatures. Finally, their behavior is intrinsically easier to model than non-solid armatures. 1 They can, however, develop into non-solid (hybrid) type armatures through the wear process at the sliding interface.
234
I I The wear process associated with the sliding couple limits the multi-shot potential
3
of the railgun. Two processes have been proposed as major contributors to the wear process. They are pulsed joule heating and arc erosion. Abrasive wear, adhesive wear, fatigue wear, and erosion are mechanisms that manifest themselves as a result of these two processes. Pulsed joule heating is sometimes referred to as the skin effect because a fast
3
rising current pulse travels in a region close to the surface of a conductor. The skin effect has been shown to result in surface melting and resolidification. 2 Arc erosion results from electrical arcing. The arc strikes from the rail to a conductor causing material vaporization
3 3 I
3
and subsequent condensation. It has been reported that the function between arc distance an inverse relationship. 3 and severity is
ARMATURE CQNSTIDERATIONS
The rise in temperature due to thermal loading has been believed to limit the armature performance. Therefore, parameters such as density, heat capacity, and resistivity have been grouped to form what is called an action constant. The action constant, g, is a measure of the loading that an armature can sustain without melting. Barber defines the
I
action constant as
To+ATpCp g=JT - dt
I and goes into detail deriving an equation for minimum mass. 4 Although thermal loading is
I
3
3
by far the most critical, Barber also presents data for -mperature rise associated with velocity, current density, and pressure. Barber's underlying assumption for minimum '235
U 3 3
mass is that the current carrying cross section carries a uniform cuirent density. The velocity skin effect negates this assumption. The velocity skin effect states that all of the current will flow in the rear of the armature when the armature is moving'. There are two options to control the velocity skin effect. The first is varying the material resistivity and the second is varying the size of the conductors. 5 Both in essence make the rear of the
I
6 armature more resistive and force the current to flow forward of the rear of the armature.
EXPERIMENTAL SET-UP
Preliminary studies were conducted using a one meter long, 12.5 mm x 12.5 mm square
I
3
closed bore railgun (Figure 1). The railgun was connected to 10 discharge capacitance banks with a total available energy of 600 Id. These banks can be charged io the desired energy level and fired in a predetermined sequence. The firing sequence is based upon a calculated desirable acceleration profile. The rail material was flame sprayed molybdenum on ETP copper. The molybdenum was used because it was believed to possess desirable
I
physical properties and resistance to arc erosion. 7 The armature sliding against the rails was composed of chevrons of 3 different materials (Figure 2). The most resistive material was at the rear of the armature. The least resistive was at the front. The three armatures
*tested
3
I
3
were as follows: Armature Materials
Resistances (microhin-cm)
Cu-Mo-Ti(6-4) Cu-Mo-Ti(CP) Cu-AI-Ti(CP)
1.67 - 5.20 - 170.0 1.67 - 5.20 - 24.1 1.67 - 2.65 - 24.1
The "graded resistaoce" design of the armatures was used to force a more uniform current distribution in the armature during flight. The design used an interference fit of up to 0.508
236
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I I I I U U I I I U I I I Figure 2: Armature Design.
I I I 238
I mm to ensure contact of the armature to the rails. For solid armatures contact to the rails is
I
of primary importance. If the interference fit is minimal and the armature loses contact with the rails, severe arc damage results. If the interference fit is too large the armature may weld to the rails. The design also incorporates the "leaves" of the armature shaped in the form of a "C". An armature with square edges at the rail/armature interface will cause the current to be concentrated there, resulting in severe heating and possible local material
I
removal. 8
I
RESULTS
The experiments were conducted using a bolted armature and flame sprayed
I
molybdenum on copper rails. The molybdenum coating was 1.016 mm thick. The first armature, Cu-Mo-Ti(6-4) had a 0.508 mm interference fit. The projectile broke apart inside of the gun and the recovered pieces were studied. The following can be observed. The copper leaves show abrasive wear tracks suggesting thermal softening of the material. (Figure 3, top) The molybdenum leaves show extensive damage. Melted and resolidified droplets of molybdenum appear on the surface of the trailing titanium leaves. (Figure 3, bottom) There is some evidence of melting of the molybdenum leaves, possibly due to a
3
premotion microwelding of the molybdenum leaves to the molybdenum rail coating (Figure 3, middle). The titanium leaves do not show any evidence of damage. Heavy debris accumulation suggests that these leaves acted as a mechanical scavenging device. The bolt connecting the leaves together was the only part of the armature to become accelerated out of the bore of the gun. It was hypothesized that the titanium was too resistive and the current never flowed in it. Because there was not current flow in them they were not being accelerated and resulted in the armature being pulled apart. 239
IOP
Fiue3IerSrae fCifr voydmnad iaiin
24
I The second armature, Cu-Mo-Ti(CP), was used as an improvement over the first.
3
By changing only the type of titanium (to a lower resistivity) and reducing thi interference fit to 0.254 mm a comparison could be drawn. The second armature behaved identically as the first. The bolt was again accelerated out of the railgun, leaving behind the material that it held together.
I
The last experiment of the "graded resistance" armatures was one that changed the strength of the bolt and substituted aluminum for molybdenum. The bolt was changed because the armature was believed to be breaking apart due to the failure of the bolt and the molybdenum was changed to alleviate the microwelding problems. The interference was reduced to 0.01905 mm. The minimal interference fit was used to reduce inhibition to movement, but led to the melting of the armature due to severe arc damage. Debris inside
I
of the bore was recovered and examined. Also, armature fragments were recovered outside of the gun and observed.
3
Some of the features common to all of the "graded resistance" armatures is severe wear of the molybdenum coating corresponding to the at rest position of the armature, failure of the bolt due to thermal loading, and unacceptable oxidation formation of the
I
3 I I
molybdenum rail coating exposed to the atmosphere. It was decided that the best way to tackle these problems was to use an armature that did not need to be bolted together and to eliminate the molybdenum coating and use plain ETP copper rails.
ARMATURE DEVELOPMENT
Due to the findings of the previously described experiments, two items become
3 3
important with respect to the design of the armature. First, the armature must remain in contact with the bore of the gun. Second, the armature should not rely on the mechanical 241
I bonding of the different materials together. Another concept that still held favor was the idea of graded resistance. Two such graded resistance armatures are under development at The University of Texas. The first is the "contact assured" armature that varies the number of conductors to grade the resistance and the second is a metal-ceramic composite composition of which is continuously varied to achieve a graded resistance. The contact-assured armature isproduced by inter-leaving a tungsten mesh between copper foils. The schematic geometry is shown in Figure 4. As shown, the interfiber
spacing is dictated by the weave of the mesh material. The inter-weave spacing can be varied to any distance needed. Varying the inter-weave spacing allows the resistance to vary continuously from the back to the front of the armature. The tungsten mesh is a 60 x 60 mesh with a fiber diameter of 0.0508 mm. This stack up of materials has been consolidated using a homopolar generator in conjunction with a consolidation press fixture. The homopolar generator passes a large current through the material. The copper softens and flows around the tungsten fibers forming a solid piece of material. The material system of copper and tungsten was chosen because of the absence of copper-tungsten intermetallics and the favorable physical properties of copper. After the material is consolidated, the faces in contact with the rails are etched to reveal the fibers. A five
minute nitric acid etch was used to remove the copper and expose the fibers. Varying the etchant time allows more or less fiber to be exposed. 'The exposed fibers are the mechanism assuring the contact to the rail material. Figure 5 shows such an exposed sliding surface.
The second solid armature under development is a combination of metal and ceramic. The metal is aluminum and the ceramic is alumina. The original form of the materials is a powder. By mixing different percentages of alumina into aluminum and layering the blended powders, a solid piece of material can be processed. Again the
I I
242
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0001 0
00
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00
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I0
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00
0
0
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I
0 0
0
0
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0
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I
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Figure 4: Foil and Mesh Combination
I I I I
243
L
I. i
I I
I I Figure 5: Copper and Tungsten Consolidation
i
I I I I I I
244
mechanism of consolidation is the homopolar generator and consolidation press fixture.
-
The combination of a low and high melting temperature material allows the aluminum to soften and flow without affecting the ceramic. This solid piece of material would have to be machined into a leaf design incorporating an interference fit. Preliminary proof-ofprinciple experiments have indicated that both these approaches have the potential to produce graded resistance composite solid armatures.
I I
Acknowledgments This research was funded partially by the following contracts: DARPA/NADC #
N62269-85-C-0222,ARDEC/SETA # DAAA21-86-C-0259, ARDEC # DAAA21-86-C-
I
0251, and DARPA/ARO # DAAL-03-87-K-0073. The authors would also like to thank Dr. M. Schmerling for his assistance with the microscopy and the entire GEM technical
staff for their help in conducting the experiments.
I I I I I I I i
245
I
1.
Barber, J.P., "The Role of Armature Losses in Railgun Performance," Presented at EM Armature Workshop, Elgin AFB FL, 24-26 June 1986.
2.
Persad, C., Peterson, D.R., "High Energy Rate Surface Modifications of Surface Layers of Conductors," Electromagnetic Launch Technology, 3rd Proceeding, April 1986, Institute of Electrical and Electronics Engineering, Inc., New York. Pp. 203205. IEEE Transactions on Magnetics, Vol. MAG-23. November 1986.
3.
Bedford, A. J., "Rail Damage in a Small Calibre Railgun," IEEE Transactions on Magnetics, Vol. Mag-20, No. 2, March 1984.
4.
Barber, J. P., "Metal Armature Critical Issues," Presented at EM Armature Workshop, Elgin AFB FL, 24-26 June 1986.
5. Condit, B., and Stachowski, R., "Velocity Skin Effect: Solid and Plasma Armatures," Presented at EM Armature Workshop, Elgin AFB FL, 24-26 June 1986. 6.
Long, G. C., "Railgun Current Density Measurements," 3rd Symposium Proceeding, April 1986, Institute of E!ectrical and Electronics Engineering, Inc., New York. IEEE Transactions on Magnetics, pp. 1597-1602. November 1986.
7.
Donaldson, A. L., et al., "Electrode Erosion in High Current, High Energy Transient Arcs," 3rd Symposium Proceedings, April 1986, Institute of Electrical and Electronics Engineering, Inc., New York, IEEE Transactions on Magnetics, pp. 1441-1447, November 1986.
8.
Long, G. C., Ph.D. Dissertation, The University of Texas at Austin, 1987.
I I
I U I I I I I
246
TASACfONS ON MAGNETICS. VOL 25. NO. 1.JANUARY 19
COM STE SOLID ARMAIRE CONSOUDATK N BY PUSE POWER 'R(XTFSNO: J ICATKhi IN .ML 'IMIlNt W;Y -AR GENERAlU AIM A NOVEL WHOM C. Pasad, D. R. Pewson, ad R. C. Zowaka Cenu for EtcIumehanics, The Umvasity of Texsm a Austin, Austin. TX 78758. armatures from powder-based metal-ceramic laminar microcomposites, and from metal-metal composites were candidates.
ARATRACT i e of the desirable characteristics for the armatures used in railguns appear 3olid feasible through the use of composite materials. Graded electrical resistance, and ssured sliding contact are among these characteristics. Metal-metal, metal-ceramic, &ad metal-polymer composites are generic types solid armature materials. of potential Copper-tungsten and aluminum-alumina are eXample composite systems, of the metal-metal Snd metal-ceramic types respectively, which have been employed in the demonstration of the versatility of the Homopolar Generator as a materials processing tool for composites. powder metallurgy and laminate bonding approaches have been utilized. In conventional processing, energy deposition is prolonged and the bulk heating which results often destroys the structural and chemical integrity of the starting materials. Composite solid armature have been consolidated with materials exposure. temperature high subsecond Densification in the solid state proceeds by a warm/hot forging mechanism and fully dense composites are obtained by a combined application of pressure and a controlled energy input,
Such composites have been produced by a novel employs a experimental approach which homopolar generator in a pulse-powered materials consolidation system. The feasibility of this approach has been demonstrated for a wide variety of engineering materials [2,3]. A relevant composite system is the binderless copper-graphite composite brush material, which shows application potential for high-speed, high-current duty (4,5]. RXPRRTMENTALPROCEDURE
Thp Conolidation Process shows the schematic tooling Figure 1 consolidation process. configuration of the A 10 MJ homopolar generator dedicated to industrial applications was the pulsed power source used, in conjunction with a modified 100 The ton vertical axis hydraulic press. operating characteristics of the HPG as a powder processing tool are described elsewhere (2].
ITRODUTIN Performance evaluation of railguns depends on availability of suitable solid armatures. Rail/ armature combinations which promote efficient energy transfer also minimize bore damage. Less flexibility exists in the choice of rail materials than in armatures. In the search for
optimum performance, it is therefore logical to
focus on alternative armature designs, and on
methods
for their processing and fabrication.
A broad-based, integrated approach to solid armature development has evolved over the last for Center the At years. several Electromechanics at the University of Texas (CEM-UT), predictive electrothermal models have been developed and successfully applied to actual armature performance. The models have confirmed the desirability of conductivity grading for improving the current distribution within these ultrahigh speed, variably-loaded sliding contacts. The requirements can probably best be met by composite armatures. The underlying concepts have been tested in a series of preliminary experiments [I1.These experiments employed a macro-composite solid armature constructed of leaves of materials with different electrical conductivities. Subsequently, fabrication techniques for a more controlled and customized grading of electrical conductivity were proposed. Solid
I
For the powder-based metal-ceramic composite consolidation, powders were mechanically mixed in batches, with ceramic mass fractions of 1% to 10%. Predetermined amounts of each batch were loaded into the cavity of the die to produce an eleven layer stack with graded composition (0% to 10% ceramic mass fraction). The insulated die cavity was lined with a dense alumina tube with an inner diameter of 50
Mm. The laminar copper-tungsten composite was constructed from copper foils and tungsten mesh. These materials were stacked such that the controlled variation in spacing between the exposed tungsten fibers in the consolidated composite would yield a graded electrical resistivity. The concept is illustrated in Figure 2. The dies and starting materials were initially at ambient temperature. Consolidation was For performed in laboratory air. consolidations performed at pressures : 300 MPa ( 45 ksi ), copper electrode/plungers were used for energy/pressure transfer to the precursors. For pressures > 300 M~a ( 45 ksi AISI 416 stainless steel electrodes were employed. After compacting the precursors to a desired pressure, the large current pulse was used to heat and consolidate the composite material. Following the discharge, the pressure was
247
I-
------------,x
Figure I.Schematic of the arrangement of the HPG powered consolidation apparatus. maintained for a holding period, permitting conductive heat transfer through the plungers to massive copper platens. The die was then was unloaded ejected. and the consolidated compact
The voltage drop across the compact and the magnitude of the current were recorded during the pulse. The product of these values was used to obtain the power during the pulse. The was numerically integrated to produce power the total energy input to the specimen, Division of the total energy input by the mass consolidated yields the specific energy input (SEI) in units of J/g. A set of curves for a copper-tungsten consolidation are shown in Figure 3.
Microstructure and density measurements were used in the preliminary evaluation of these composites. The consolidated composite materials were sectioned in the radial direction, metallographically prepared, and then examined by scanning electron microscopy.
Anat.i
nrBcursors
One potential difficulty with these materials in high speed sliding duty is the relatively large scale of the constituents. The apparent sliding contact behavior, fn which 5-20 point contacts actively carry current at any instant [6), suggests that more finely dispersed structures may produce better performance. Furthermore, in the refractory-metal skeleton systems, such as copper-infiltrated tungsten, where the vaporization of the infiltrant effectively controls wear, the need is for many micro-islands of tungsten each surrounded by copper. These difficulties may be resolved by the adoption of novel materials engineerin; approaches. For the refractory-metal-based systems, the ultrafine filamentary composites described by Bevk [7] is an attractive technical alternative. These materials are
RESULTS AND DISCUSSION
I
MinrnmtIrntIure
Highly dense microstructures were observed throughout the cross-sections of the consolidated composite materials examined by SEM as shown in Figure 4 for the copper-tungsten composite. Excellent conformity between coppc- and tungsten indicates the potential for a highly conductive interface. An etched cross section is shown at three magnifications in Figure 4 (a), (b), and (c). Figure 4 (c) also indicates the retention of the interfiber spacing of the starting material.For both material systems, densities greater than 95% of theoretical were achieved in the preliminary experiments.
3
248
Figure 2. Schematic showing the interleavi concept applied to the copper foil/ tungst weave composite starting materials.
oracterized by dense and uniform dispersions
S
*
0to 1010/cm2) of fine filaments (5 to 100 nm *j ck) . In this approach a powder metallurgy
COSLDTN
1 CONSOLIDATION 2
Soduct may be used as the starting material. is material is deformation-processed by .lling, swaging, drawing, or extrusion to very rge strains until the in-situ formed •tilaments are sufficiently small.
SITE
Lt 1.0
Another approach to microstructural refinement is to derive the metal powder blend from the reduction of a complex metal oxide precursor. copper tungstates and copper suitable molybdates for this purpose have been identified. The challenge is then to distributed, finely these consolidate homogeneous powder blends into useful forms without structural coarsening. This is a task to which the high-energy high-rate consolidation processing may be well matched.
0
0.o 400
2 a
300 200-
C
100
U
jn the case of the aluminum-alumina system, two I alternative approaches may be taken to distribution of the alumina in the metal matrix. One is by heat treatment of air-atomized powder particles after compaction to intermediate densities (-85%). This leads t spheroidization of the alumina in a continuous aluminum matrix, which can then be consolidated to full density. The second approach is by welding and mechanical alloying. The continuous assemblage, fracturing of an aluminum powder of the and the rapid oxidation freshly-fractured surfaces leads to a fine and uniform dispersion of the alumina (8]. These composite powder particles could then be consolidation processed into desired shapes.
0. ; 120,000. . I 8 , w a 40,000 0 0.0
0.4
0.5
1.2
1.6
TIME (Secondal
Figure 3. Current, power and total consolidation energy waveforms applied to the copper foil/ tungsten weave composite starting
Metal-polymer copngsites The widespread use of polycarbonate cubes as the accelerated load in plasma-driven railguns indicates the survival potential of polmer-on-metal sliding couples in the hypervelocity regime. Recent advances in conducting polymers raise the question of the potential benefit of the incorporation of such materials into solid armatures. It has been reported that on a mass basis a new, iodine-doped polyacetylene polymer is as conductive as copper. More recently, Lockheed has developed the polyanilines which in large scale production may be dollar-a-pound material. These materials are claimed to be less difficult to process than other conducting polymers and they do not ignite upon exposure to air (9]. The
potential exists
to
incorporate these
advances in polymer technology into the development of a composite solid armature design which matches the high-speed, high-current duty demanded of these components. One approach may be the use of a tubular skeleton made of a fine tungsten powder, infiltrated with a suitable conducting polymer and fitted over an aluminum or aluminum-alumina composite core. Such an arrangement takes advantage of the tribological and conducting properties of the polymer, the high melting temperature of the tungsten, and the low mass of the core. Other interesting possibilities exist for grading the mechanical, electrical and tribological properties in the radial and longitudinal directicn.
249
1
MM
300 p
100 PM
Figure 4. Scanning electron micrographs at (a) low (b) intermediate and (c) high magnification of an etched cross section of a copper foil/ tungsten weave based planar composite.
I 7.
Bevk,
"Ultrafine
Filamentary
Some of the desirable characteristics for the
Composites",
solid armatures used in railguns appear feasible through the use of composite materials. Graded electrical resistance, and assured sliding contact are among these characteristics. Metal-metal, metal-ceramic, and metal-polymer composites are generic types of potential solid armature materials, Copper-tungsten and aluminum-alumina are example composite systems, of the metal-metal and metal-ceramic types respectively, which have been employed in the demonstration of the versatility of the Homopolar Generator as a materials processing tool for composites. materialsGMENT pReport, ACKNOWLEDGMENT
Materials Science, 13, 1983 edited by R. A. Huggins, R. H. Bube, and D. A. Vermilyea (Annual Reviews Inc., Palo Alto, CA) pp. 319-338).
This research received early support from SETA Contract DAAA- 86-C-0259. The latter stages were supported by DARPA/ARO Contract DAAL 0387-K-0073. The authors thank Ted Aanstoos and Jim Allen for their assistance with the processing. Chris Lund and M-J. Wang assisted with the microstructural evaluation. This work is part of a collaborative effort between the Center for Materials Science and Engineering and the Center for Electromechanics at The University of Texas at Austin. We thank them for their cooveration.
2. C.J. Lund, C. Persad, Z. Eliezer, D. R. " Peterson, and J. Composite Solid Armatures forHahne, Railguns", Paper #13, Proceedings of the 3rd. Int. Current Collector Conference, Austin, TX, Nov. c1987. I
j.
2. H. L. Marcus, D. L. Bourell, Z . Eliezer, C. Persad, and W. F. Weldon, Journal "High-Energy High-Rate Processing", of Metals, December 1987, pp 6-10. 3. H. L. Marcus, C. Persad, Z. Eliezer, D. L. Bourell, and W. F. Weldon, High-Energy High-Rate Processing of Composites", Proceedings of the AIME Northeast Meeting, Hoboken, NJ, May 1987. 4. K. C. Owen, M. J. Wang, C. Persad and Z.Eliezer, "Preparation and Tribological Evaluation of Copper-Graphite Composites by High-Energy/High-Rate Consolidation", Wear 20, 1987, pp. 117-121. 5. M. J. Wang, C. Persad, Z. Eliezer, and W.F.Weldon, "High-Energy / High-Rate Consolidation of Copper-Graphite Composite Brushes for High Speed, High Current Applications", Paper #20, Proceedings of the 3rd.Int. Current Collector Conference, Austin, TX, Nov. 1987. 6. I. R. McNab, " Pulsed High Power Brush Research II. Interpretation of Experimental Data", Energy PulsA
in High Power-High Production and
AppiJcati n, ed. E.K. Press, Inall, Australian National University (Canberra) 1978. pp 201-215.
250
in
Annual
Reviews
of
and Properties of Aluminum Alloys Produces by Mechanical Alloying: Powder Processing and Resultant Powder Structures", Metall. Trans., 12A, May 1981, pp. 813-824. 9. T.S. Kuan,'Conducting Plastics Sell for $1 a Pound', in SUPERCONDUCTIVITY Flash Vol. 1, No. 9, Jan. 11, 1988. p.8.
. WEAR OF CONDUCTORS IN RAILGUNS: METALLURGICAL ASPECTS C. Persad, C.J. Lund, Z.Eliezer, Center for Materials Science and Engineering, University of Texas, Austin, TX 78712 and D.R. Peterson, J.Hahne, R.Zowarka, Center for Electromechanics, University of Texas, Austin, TX 78758
interface between dissimilar material pairs; material exchange between the armature and the rail; chemical reactions between the
ABSTRACT
The wear of metallic rail conductors influences the performance of electromagnetic railguns inwhich initially solid armatures (sliding electrical contacts) are accelerated to high velocities. Vapor and f'me liquid droplets formed at the armature/rail interface reduce the multishot operating efficiency.The experimental focus in this study was the analysis of Aar debris generatedby an aluminum alloy
armature sliding on copper rails. The experimental parameters were set to maximize armature velocity. Preliminary rail and armature wear observations were documented as part of a database for solid armature development. A variety of debris collection and characterization approaches were used. Both in-bore debris and muzzle- exit debris were collected for size and chemical analysis. Debris constituents were copper rich and the compound Cu9 AI4 was observed in the splat-quenched muzzle debris. The approach and method are detailed to indicate the potential of wear studies in defining the activity at the sliding interface. Some implications of rail surface roughening on zero wear measurements are also discussed and the adoption of wear mechanism maps is proposed. ITO
I
Wear processes which occur in railguns are characteristic of sliding electrical contacts [1,2]. These processes are complex and appear to be system specific. From a metallurgical viewpoint it is interesting to view a high-speed solid armature as a vehicle for surface modification of rails. Such an approach ispotentially important for predicting performance in the multi-shot operating environment where successive changes in the surface state and chemistry of the rails can drastically alter the mode in which the current isdistributed in the near-surface regions. This in turn changes the coupling into the solid sliding armature. When a particular armature/rail couple is subjected to an electro-thermo-mechanical loading spectrum the metallurgical response isonly qualitatively predictable based upon enlightened empiricism. The difficulties associated with more precise predictions are evident when the range of contributing phenomena is considered. Important metallurgical effects include alloy formation due to elevated temperatures developed at the sliding
251
environment and the freshly re-exposed rail surfaces; and the embedding of debris produced in the forward section of the solid armature and pushed into the softened rail surface by the high pressures in the rear section. Muzzle arc debris produced at armature exit further complicates the metallurgical analysis. The temperature and pressure changes at the sliding interface are particularly important to armature performance. Pressure effects
originate in mechanical preload and in magnetic pressure. The temperature effects have been analyzed by Kuhlmann-Wilsdorf [I]. Pressure effects have been included in the analysis by Ross et. al. 171. The flash temperature solutions proposed by Kuhlmann-Wilsdorf allow heat evolution at contact spots. Joule heating is assumed to be velocity-independent, while friction heat is proportional to the relative velocity. Contributions to Joule heating include film resistance heating and constriction resistance heating. Qualitative accomodationofthecontinuousquenchingofthearmature sliderby cold rails make these temperature predictions upper-bound solutions. The dynamic friction force calculations of Ross et al. are translated into friction coefficients (ps) through constant normal contact force Copper -on -copper sliding at up to 0.20 km/s yields a nearly linear drop in p from 0.40 to 0.10 [2]. These coefficients are low at velocities > 0.10 km/s. Uncertainties inthe value of the inductance per unit length, L' (assumed constant), and in contact force (alo assumed constant) are suggested reasons for the low values. These reviews indicate that existing quantitative predictive models are still incomplete. A single parameter predictor such as interface temperature remains attractive. The difficulty lies inascertaining the contributions of each of the multiple and coupled heat-generating sources. Potentially more difficult is the prediction of the conditions where impact-generated stresses and consequent localized energy absorbtion can push the temperatures sufficiently high to cause solid-to-hybrid conversion of armatures. It is apparent that bore smoothness, bore straightness, and bore stiffness concepts being incorporated into the present generation of EM guns may contribute to extended bore life by reducing or eliminating some of the sources of impact-generated stresses.
In this paper, we focus upon wear debris analysis for a monolithic • juminum alloy armature sliding between ETP (electolytic tough pitch) copper rails. The experimental parameters were set to maximize armature velocity. Preliminary rail and armature wear observations were documented as part of a database for solid armature development. A variety of debris collection and characterization approaches were used. Both in-bore debris and muzzle- exit debris were collected for size and chemical analysis. The approach and method are detailed to indicate the potential of wear studies in defining the activity at the sliding interface.Some implications of'rail surface roughening on zero wear measurements are also discussed and the adoption of wear mechanism maps-is proposed.
EXPERIMENTAL APPROACH EML System Details of the 7050 aluminum alloy 'fishbone' solid armature and of the Im long, 12.7 rm square bore railgun used for solid armature development are given by Price et. al. [3]. The current waveform for the experiment selected for detailed analysis was roughly semi-sinusoidal inshape with a peak current of 320 kA at 0.75 ms ina pulse of total length 1.9 ms. The armature was cooled in liquid nitrogen to reduce the insertion force due to its interference fit. In the interval between loading and firing, the armature returned to room temperature. The armature was sized 0.3 mm greater than the rail-to-rail dimension. The armature was accelerated to a muzzle exit velocity of 1100 mIs.
3
Debris Collection
3
Post-shot debris collection employed a I m long, 6mm diameter stainless steel tube connected to a sealed glass jar for in-bore vacuum scavenging. This technique combines a small mechanical action of the tube tip and vacuum suction to remove loose in-bore debris. muzzle-exit debrisflats was inclined stripped atfrom arc chutes. The arcThe chutes are two steel a 45thedegree angle relative to the rail surfaces, and positioned -10mn above and below the rail surfaces at the muzzle exit.
coated with a clearer transition from a thick to a thin deposit. The thickness of the deposit was greater at the breech end and became progressively thinner towards the muzzle. The process by which the material is transferred from the armature to the rail is not entirely clear. It may be 1) smeared while the armature is in the heat-softened solid state, 2) transferred from a liquid interlayer to the rails ina mode qualitatively similar to planar flow casting or 3) spray deposited as a fine debris driven off a transient molten layer and entrained in a debris plume which spreads the debris well away from its source. It appears likely that as the initially solid armature progresses down the bore and the sliding interface temperature rises, the dominant process changes from 1)to 2) to 3). However, non-uniform longitudinal and lateral temperature distributions make it possible for these processes to operate simultaneously, as evidenced by the streaking of the transferred material. These observations indicate that the degree of wear of a solid armature produced by armature/rail interaction is indeed a function of the energy dissipated during the residence time of the armature at a particular location. This confirms the desirability of measures taken to avoid armatures starting from rest in a railgun bore. Typically the motivation is to reduce rail damage. Pre-acceleration might produce the added benefit of reduced wear of the solid armature. Debris Analysis. SEM analysis shows the in-bore debris dimensions to range from tens of microns to submicron. Figure 1 shows a typical high magnification view of the extracted, loose, in-bore debrif mounted
Debris Analysis The scanning electron microscopy examination was conducted in a JEOL 35M instrument operated at 25 kV. Chemical analysis was performed by energy dispersive spectroscopy in the SEM. This provided the size and overall chemistry information. X-ray diffraction data were obtained on a Phillips diffractometer. This allowed confirmation of the crystal structure of the phases present.
3
RESULTS AND DISCUSSION
Wear Related Observations
3The
condition of railgun exterior (i.e. the steel structural envelope) was documented as was the state of gun bore surfaces viewed by a borescope from the breech and from the muzzle ends. The pertinent observations are summarized below. Exterior: The arc chutes were covered with a splat-quenched deposit which had a metallic copper-aluminum appearance. The top chute which was located slightly closer to the bore had a larger and thicker deposit. The rounded rail edges at the muzzle exit appeared to be heavily eroded. Th.re was no evidence of any black powdery soot normally associated with carbonaceous debris production.
Gun Bore: The general appearance of the bore was dominated by an aluminum-colored coating, with occasional gold-colored streaks. *,The deposit was not uniformly distributed on the top and bottom rails or along the gun length. The lower rail was more heavily
3
252
Figure 1.Scanning electron photomicrograph of loose in-bore debiS.
3
iocolloiualgraphte. Thlarge paclsappar tobe agglomeCratos 6f small particles. We focus upon the micron-sized particles which
1
e
V.
0 counts
I
D..r.
Coop- 3
5
Elapsed-
-ALCO
(a)
Vpcar to be the building blocks of the debris. This particle
slorphology leads us to hypothesize that these particles were splerodized inflight inthe bore. They were deposited onto a more tightly adhered and already oxidized wear layer which made them easier to vacuum scavenge out of the bore. Two different debris chemistries are evident in the EDS scans given inFigures 2a &b.In "iure 2a, Cu-Al- Zn rich debris is evident. The source of the zinc isthe 7050 alloy chemistry (wei ht percent 6.2 Zn, 2.25 Mg, 2.3 Cul, 0. 12 Zr. Bal. Al). In Figure 1b, taken from the "gold-colored" debris, the Zn is absent.
..u
Uv
-
~~.-
The muzzle exil debris is 80 Vin to 100 pim thick covering areas of
200-400 mm 1 .The spectrum of elemental constituents is grphically detailed with micrograph inset inFigure 3.The X-ray diffraction studies focused on the muzzle exit debris. Copper X-ray
14-
8.120
_ _ _ _I8 o
Rar'.
10,230
1ev
9___2__
I___Integral__________
diffraction lines were present in each of the two-theta diffraction
spectra. The line broadening observed was attributed to the residual
siains associated either with a supersaturation of alnminum or with rapid splat quenching. In addition to the copper-rich matrix, the intermetallic compound Cu9AI4 was found to be present. Cu9AI4 had been previously observed in a railgun deposition environment by DeLuca in his studies of the quenching of exploded aluminumIC foil onto a copper substrate [4].
U%
I
.khL'ODBRDEIS2 280 countsQ4 DisciS I2
V,%
Dn
66 sect
C£10e
9
1(b) IC
properties of the Sliding Pair AIICu
UBore
wear analysis requires that the physical properties of the
armature/rail material pair be understood, since the severe wear observed appears to be dominated by large-scale material transfer.AtC This section reviews the equilibrium phase diagram, and examines the influence of alloy chemistry on fluidity and resistivity.
I 3Vert-
C _
.d
''
4-
0.328
Range-
54 -n.rI 10.239
10.230 keV
Energy Dispersive Spec troscopy (EDS) Plots indicating the Figure 2.
MUZZE DEOSITchemistry of the loose in-bore debris (a)Cu-Al-Zn rich (b)Cu-Al rich.
2000 counts
Disp- 1
Elapsed-
60 secs
K..
...... U -: 1::*.
4-
::*a::"...'--
0.120
Range=
10.230 keV
10.230
Figure 3. Energy Dispersive Spec troscopy (EDS) Plot indicating the chemistry of the muzzle exit debris is Cu-Al rich. The appearance of the debris is shown inthe inset photomnicrograph.
253
percent copper) fluidity increases significantly. The eutectic fluidity
Pha Diagram
The equilibrium binary phase diagram for the system Al-Cu is shown in Figure 4, taken from reference [5] in which the characteristics of this system are detailed. Pertinent to this discussion is the described multiplicity of complex solid state
reactions. Here we focus on the source of the Cu 9AI4 . Cu9AI4 , based on the gamma-brass structure, is the gamma- I phase which exists over the compositional range of 62.5 to 69 at. % Cu. Its appears to be the stable high-temperature 8 phase which source melts at 1322K. A two phase [(Cu) + B] field exists between the eutectic temperature and the eutectoid reaction S ->gamma- I + (Cu) at 850K. The gamma-l phase has also been observed in diffusion in the prehae of The 773K0K5). asee couplesannealedat couples annealed at 773K [5]. The presence of Cu9 Ai 4 iof splat-quenched muzzle debris is a clear indication of an alloy liquid phase formation. From the phase diagram it may also be observed that the eutectic AI-Cu alloy melts at 112K lower than Al. The formation of this eutectic may place an effective limit for the temperature increase AT in the solid-to-solid sliding regime, the
liquid thresholdbeing its formation rather than the melting
of Al.
For aluminum-copper binary alloys, fluidity decreases rapidly as the copper is added, reaching a minimum at -10 mass percent copper. As the alloy composition approaches the eutectic (33 mass .,r:,,t
I.,c,€
Resistivity The room-temperature resistivity of aluminum as a function of mas of copper added is shown in Figure 5, which is a plot of data from reference [7] . Resistivity is observed to rise steadily with increasing copper content, and to increase significantly in the range 70 to 95 mass %Cu. If the interface voltage drop is to be correctly predicted for an armature which develops a liquid surface, then the resistivity changes due to alloy formation must also be taken into account. Furthermore, if such alloys are formed and are a residual deposit on the rail surface, then the implications for advantageous current redistribution also need analysis.
The potential roles of electrical and frictional energy dissipation at the sliding interface have been mentioned. The discussion now includes the least well-understood wear phenomenon: gouging due to impact energy dissipation. In uncoated copper rails, postfiring observations reveal gouges appearing inclusters of up to six small I mm to 3 mm in diameter. Their appearance is easily from arc damage, which, in copper rails, takes the form of small, densely-packed craters lOum to lOOum in diameter. Clusters of gouges appear on alternate sides of the bore in a
*pits,
Copir
phenomenon at the armature/rail interface.
Impact Energy Dissipation.-Gouging
Fluidity
I
is higher than that of pure aluminum [6]. Fluidity and chemistry in the AI-Cu system are important in explaining the the smearing of alloy debris in the bore and the possibilities for development of a superheated transient thin liquid film layer as a friction reducing
-distinguishable
LL L
distinctive repeating pattern. The disiance between successive clusters increases with increasing armature velocity, suggesting the presence of a secondary transverse perturbaion of constant frequency superimposed upon the primary accelerating force.
The phenomena of gouging due to intermittent grazing impact at
velocities > 1 km/s has been reported previously [8]. It was observed that material ejected from the gouge was deposited in the form of a metallic spray forward of the gouge pits. In early railguft research [9], it was reported that for copper sliding on copper,. gouging begins at 0.6 km/s and increases in severity with velocity., Indium, rhodium, and gold plating of copper rails were each found
-
-.
Ile
[
- ..
-
J
)_--
ifound to eliminate the gouging phenomenon. The absence of a
, .. -fundamental Ia Cu
Weight Percent Copper
A)
to marginally reduce gouging. In our research, the use of a 0.5 nll thick, flame-sprayed molybdenum coating on copper rails was
Figure 4. The equilibrium binary phase diagram for the system Al-Cu 40__
to a variety of experimental conditions. These include: the bore finish and fit of the armature and-the general integrity of the
structural envelope. Straightness, smoothness and stiffness art desirable attributes and may be crucial to delaying the onset of
~gouging. INFLUENCE OF Cu ON
understanding of gouging prevents a rationalization of this finding. At the present time quantitative data on bore damage is scarce. Qualitatively, the degree of bore damage appears sensitive
RT RESISTIVITY OF A] at
Crack Formation
30
Wear processes are likely to be influenced by large fluctuatii. stress conditions as a result of both magnetic pressure and therim' changes. Pressures on the order of 40 ksi [10] , and therl
Ci
gradients of 1000K over 200 microns are not uncommon [11].
conditions have led to surface microcracking of the rails am
20 -These
insulator side walls [12]. The continued effect can result in surface
failure and break-up into wear particles. ix 10
In pulsed joule heating experiments at I to 10 MW/mm 2, metallic conductors were found to undergo a variety of surface transformations [I ].The effect of pulsed joule heating alone, iD the absence of arc tracking or high speed sliding contact, WO sufficient to produce surface melting and microcracking. It appe"
N'
0
_L
0
I,.
20
_
_,__
40
60
I__
80
100
Copper (wy %) Figure 5.The room-temperature resistivity of aluminum as a function of mass of copper added {data from reference (7)).
254
that these effects were aggravated by the presence of non-mellrA second-phase inclusions in commercial .purity materials. Su4;, inclusions presumably cause constrictions in surface current flow and may contribute to localized surface melting. The developmeo' of directionally resolidified grain structure [13,14] may aI,5 influence crack resistance.
materials, etc., a eries of sub-plots will be required. Models which The Zero Wear' Artifact geports of 'zero wear' in railgun bores in which several MJ )i energy have been dissipated have become disturbingly common. The abundant production of wear debris, whether in spectacular muzzle exit debris or the more common 'railgunner's soot', suggests a need for re-interpretation of these results. Two invisible' sources of wear debris have come to our attention: the first is the micro-roughening due to liquid droplet impact or to microparticle or ion clusters being plucked from the rail surfaces; the second is the wear material removal which accompanies the rounding of the rail corners, where both theory and experiment indicate that the current density is at its maximum. These wear phemonena may not change the rail-to-rail dimensions but do alter the surface roughness. These phenomena have been non-destructively checked through the use of surface replication
predict energy dissipation are currently being developed and await experimental verification. When these are implemented, wear mechanism maps can be used to graphically define the transitions at the sliding interface and to point the directions for wear control in railguns. ACKNOWLEDGMENT This research received early support from SETA Contract DAAA86-C-0259. The latter stages were supported by DARPA/ARO Contract DAAL 0387-K-0073. The authors thank Drs. M-J. Wang and M. Schmerling for their assistance with the microstructural evaluation. This work is part of a collaborative effort between the Center for Materials Science and Engineering and the Center for Electromechanics at The University of Texas at Austin. We thank them for their cooperation.
techniques used infractography [15].
The 'zero wear limit' is used in the Bayer-Ku sliding wear theory [16]. Adopting a similar definition for railgun rail conductors, the 'zero wear limit' will be the point in the life of a railgun when its surface roughness median has been depressed to half the depth of the initial peak-to-valley finish, as illustrated in Figure 6 adapted from Engel [17]. Such a definition allows accomodation of the fact that microscopic changes in asperity dimensions precede the development of distinguishable wear. This 'zero wear limit' will be important in future analytical treatments of rail wear, since it marks the start of a measurable wear rate: a process which is crucial to
1. 2.
3.
Distributed at Proc. of the 4th. EML Conf., Austin, TX,
accurate bore life prediction. At this time, it is possible that not
many experimental EML systems have undergone sufficient cycles to produce measurable wear by these criteria. However, the potential pitfall is the behavior within the 'zero wear limit' being extrapolated.
4. 5. 6.
Wear Mechanism Maps The wear mechanism mapping approach proposed by Ashby's group [18] is potentially useful. For rail wear, appropriate axes might be mass removed vs energy dissipated. Since energy dissipation is itself a function of current, velocity, sliding pair
7. 8. 9.
10.
h
(C)
11.
12.
Measurable Wear
Zero Wear
Wear Depth
13. 13. 14
15.
h
(a) NO
Figure 6.
April, 1988. University of R. J. DeLuca, MS Thesis, 7 The Texas at Austin, May 1987, p. 30, p. 9 . J. L. Murray, Binary Alloy Phase Diagrams. T. B. Massalski (ed.), ASM, Ohio,1986, pp. 103-108. F. R. Mollard, M. C. Flemings, and E. F. Niyama, "Aluminum Fluidity in Casting," J. of Metals, (Nov. 1987), p. 3 5 . K. Schroder (ed.), Handbook of Resistivity of Binar& Alloys, CRC Press, Boca Raton, FL, 1983, p. 79 . K.F. Graff and B.B.Detlofi, " '1he Uouging Phenomenon between Metal Surfaces at High Sliding Speeds," Wear 14 (1969), pp 87-97. R.A. Marshall, "Moving Contacts in Macro-Particle Accelerators", in High Power-High Energy Pulse Production
andAplaioed. E.K. Inall, Australian National University
. h>&2
(b)
D. Kuhlmann-Wilsdorf, "Flash Temperatures due to Friction and Joule Heat at Asperity Contacts," Wear, 105 (1985) pp. 187-198. D.P. Ross, G. L. Ferrentino and F. J. Young, "Experimental Determination of the Contact Friction for an Electromagnetically Accelerated Armature," Wear, 78 (1982) pp. 189 - 200. J.H. Price, C. W. Fulcher, M. Ingram, D. Perkins, D. R. Peterson, R. C. Zowarka, Jr., and J. A. Pappas, "Design and Testing of Solid Armatures for Large-Bore Railguns",
Cycles,
N
(a)a log-log measurable and(b) weardata Zero wearschematic (c)measurable zero wearwear: forwear
255
16.
17.
18.
Press, (Canberra) ,1978. R.F.DavidsonW.A.Cook, D.A.Rabern, N.M. Schnurr, Deformations and Launcher "Predicting Bore Stresses in Railguns", IEEE Transactions on Magnetics, Vol. MAG-22, (Nov. 1986), pp. 1435-1440. C. Persad and D. R. Peterson , "High Energy Rate Modification of Surface Layers of Conductors", IEEE Trans. on Magnetics, Vol. MAG-22, (Nov. 1986), pp. 1658-1661. D. R. Peterson, D. A. Weeks, R. C. Zowarka, R. W. Cook, and W. F. Weldon, "Testing of a High Performance, Precision Railgun," IEEE Transactions on Magnetics, Vol.MAG-22, (Nov. 1986), pp. 1662-1668. A.J. Bedford, "Rail Damage in a Small Calibre Railgun," IEEE Transactions on Magnetics, Vol. Mag-20, No. 2, (March 1984), pp. 34 8 -3 5 1 R. F. Askew, B. A. Chin, B. J. Tatarchuk, J. L. Brown, and D. B. Jensen, "Rail and Insulator Erosion in Rail Guns," IEEE Trans. on Magnetics, Vol. MAG-22, (Nov. 1986), pp. 1380-1385. Metals Handbook: Fractography and Atlas of Fractographs, H. E. Boyer (ed.), 8th. ed., Vol. 9, ASM, Ohio, 1974, p.56. R. G. Bayer, A. T. Shalkey, and A. R. Wayson, " Design for Zero Wear," Machine Design, 41 (1) .(Jan.1969)pp. 142-151. Peter A. Engel, Impact Wear of Materials, Elsevier, New York, NY, 1978, p. 182.
S. C. Lim, M. F. Ashby, and J. H. Brunton, their and Transitions -Rate 'Wear Relationship to Mechanisms', Acta Metall., (June 1987), 35 (16), pp. 1343-1348.
BINDERLESS COPPER-GRAPHITE AND NANOSIZED STRUCTURES IN Cu-W-WC-C COMPOSiTES FOR ELECTROTRIBOLOGICAL APPLICATIONS: C. Persad, M. Schmerling, Z. Eliezer, Center for Materials Science and Engineering, University of Texas, Austin, TX 78712 and R. Carnes, I. Gully Center for E"cctromechanics, University of Texas, Austin, TX 78712 Abstract Sliding spee-s, current densities, and useful operating lifetimes are key parameters for electrical contacts in future high-performance pulsed power machines. For homopolar pulse generator brushes, infiltration of Pb-Sn into a copper-graphite powder composite skeleton is a conventional fabrication route. This composite fails at v>220 m/s by loss of the low melting temperature binder. Binderless copper-graphite materials operate at higher temperatures. However, speeds > 200 m/s at current densities > 2 kA/cm 2 produce interface temperatures > 5000C. Results will be presented on a recent novel materials science approach to handling of such high interface temperatures. Contact layers containing a tungsten aerogel-type skeleton structure with 100 nm-sized contact points, and with nanosized WC and colloidal graphite particles have been fabricated. Sol-gel synthesis, carbothermic reduction, large-strain deformation processing, and high-energy, high-rate powder consolidation were the main steps in the processing of these nanosized structures. The processing/microstructure/properties base for these nanosized composite structures was related to the observed high-speed, high-current, electrotribological performance.
I I I This work was supported by the Texas Advanced Research Program, Grant # 4357, 19881990, the Texas Advanced Technology Program, 1989-1991 Grant, and by DARPA under ARO Contract DAAL 0387-K-0073.
i
256
SPulseddpower Pulsed power and hypervelocity phenomena are critical emerging technologies (1). These technologies are undergirded by a variety of hardware. Among these are powerful electrical r.achines. Their powerfulness is derived in part from high operating speeds and extremely large electrical currents. Examples of such machines are homopolar pulse generators and electromagnetic launchers. Sliding electrical contacts are integral current collecting components in these machin. s. They must perfo.m under extreme conditions as depicted in Figure 1. In homopolar pulse generators, the brushes must harvest electrical currents in the megampere regime at sliding speeds > 100 m/s. In electromagnetic launchers, the solid rmatu'res must drive payloads by megampere-induced Lorentz forces to velocities in excess of 1 km/s. Schematic diagrams of each of these devices are shown in Figure 2.
High performance composite materials are being developed to meet these needs. The brushes are being fabricated as a powder metallurgy composite of copper and graphite. These constituents provide the conductivity and lubricity sought in this device (2). The solid armatures are being de.;igned as a hybrid composite component with special attention to the sliding interface (3). These "electrotribological" systems are so designated because of the combined sliding and current-carrying duty. It is well established that a number of composite materials systems exhibit the characteristics necessary for high performance sliding electrical contacts (4). When based on powder constituents, such composites provide scne of the necessary product properties by selection and assembly of materiai at the microscopic level. The relationship between the tribological behavior of such contacts and their basic materials science is the subject of a large body of studies (5-15). A key issue is the value of the maximum temperature that can be sustained by the contact at the sliding interface, while maintaining the efficient energy transfer associated with solid-to-solid contact. In these applications, the complex electrotribological environments encountered provide multiple interface heating sources, as indicated in figure 2 (a) for a homopolar generator brush. Achieving the necessary performance demands novel approaches to the design, processing, characterization and laboratory testing of suitable materials. Sliding at high speeds while canying large currents causes combined mechanical and electrical energy dissipation at the sliding interface. The contact surface becomes hot. In the case of solid armatures, the contact surface becomes hot enough to melt metallic conductors such as aluminum (16). In electromagnetic launcher terminology, the onset of melting leads to a change in electrical performance described as "transitioning". During "transitioning" a fluid interlayer (liquid, vapor or plasma) develops at the sliding interface and a solid armature then becomes a hybrid armature showing some of the characteristics of a plasma armature at the interface while the bulk remains a softened solid. Hybrid armatures perform unpredictably. It is therefore desirable to design solid armatures that remain solid during use (17). M
a
Experimental Procedures
I. Composite Brushes. The binderless copper-graphite was produced by the high-energy high-rate technique (18-20). The degree of homogeneity of the distribution of the graphite in the copper matrix depended upon the mechanical mixing that is achieved prior to loading of the powder mixture into the die. The degree of mixing depends upon the relative densities and the .... characteristics of the loose starting powders. Since the single-residence processing occurred in the solid-state, the pressure and temperature conditions within the die during the processing control the degree of densification. The copper matrix was powder-forged to high densities, while the graphite was subjected to intraparticle shear to accommodate the plastic flow of the matrix. Heating rates of 100 K/s to 500 K/s were used to rapidly bring the copper matrix to the warm processing range (0.4 to 0.6 Tm, where Tm is the melting temperature of copper = 1356K). While at temperature the composite densified under an applied stress that was in excess of the flow stress within the powder volume.
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... IH-EAT SOURCES IN A HOMOPOLAR BRUSH RESISTIVE HEATING OF STRAP PS = 12 RS = 1.8 kW
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RESISTIVE HEATING OF BRUSH PB = 12 RB = 500 W
INTERFACE HEATING FRICTIONAL LOSSES - Pf = gNv = 1.8 kW
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ELECTRICAL LOSSES - Pe = VI = 10 kW
l=5kA
V=2V RB = 20jAQ RS = 72 p(
v = 200 m/s g. = o.2 N =45 N
(a)
Electromagnetic Railgun
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PARALLEL CONDUCTING RAILS
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(b) Figure 2 - Operating environment of (a) high velocity sliding contacts (brushes) and (b)hypervelocity sliding electrical contacts (armatures). In (a) the relative magnitudes of some of the heating sources of the contact are estimated using a typical set of parameters. 259
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The measured values of density, hardness and electrical resistivity ot the b1inceriess coppergraphite composite material for brushes are listed in Table I. Table I. Properties of binderless copper-graphite composite materials for brushes. -
Density (g/cm 3 ) 6.30-6.68 94.1-99.8 (percent theoretical) Average Hardness * (Rockwell 15T) 62 (Axial), 50 (Radial) Electrical Resistivity ([t ohm-cm at 300K) 3.82 (Axial), 8.95 (Radial) On this scale a fully dense copper powder compact has an average hardness of 70.7.
3 I
Binderless Cu-Gr
1I. Composite Armatures. The composite material designed for solid armatures consists of a 100 I.m thick W/WC/C layer embedded into copper. The objective was to produce a sliding interface capable to simultaneously withstand and reduce the high interface temperature (21, 22). This W/WC/C layer was synthesized by a technique consisting of synthesis of tungstic acid gel, precursor blending with colloidal graphite, followed by spreading of a thick film of the blended precursor onto copper. The W/WC/C layer was for.aed by carbothermic reaction and embedded into copper by flat rolling followed by pulsed joule heating under pressure. W/WC/C coated copper rectangles (20 mm x 12 mm) were then cut and mounted for testing evaluation. Figure 3 shows the structure of the blended precursor thick film.
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I (a)
(b)
3
Figure 3 - SEM photomicrograph of a vacuum dried composite film synthesized from a tungstic acid/colloidal graphite blend. (a) low magnification showing the comer edges of a 100 V±m thick layer. (b) high magnification view of a fracture surface of the composite film in (a) showing the ultrafine particle microstructure. Materials Evaluation
3
Low Speed Testing. The low speed tests were conducted in air on a pin-on-disk machine. .Friction force and pin wear were monitored by a strain ring and a linear-displacement transducer respectively. An Apple II computer stored and plotted the data. The pin wear rate and friction coefficient were calculated from the plots. The applied load range was 2 to 20N, and the slid',g velocity range was 220 to 670 mm/s. High Speed Testing. The sliding electrical contact tester (SECT) is a device which allows rapid evaluation of materials and cooling methods with respect to interface power dissipation and wear rate. Continuous testing at speeds of up to 180 m/s at 6000 A is possible. Electromagnetic displacement probes monitor brush wear, and a video thermal imaging system provides a timetemperature history of the slider, sliding interface, and rotor. The tests were performed dry in air. No active cooling was used. Constant down force of 100N was maintained during the temperature excursion tests. The specimen surfaces were prepared by polishing with a final
3
260
polish using 600 grit silicon carbide paper. Prior to the test the contact and rotor surfaces were cleaned with high-purity ethyl alcohol. Post-Test Evaluation. Following tribological evaluation, the contact surfaces were examined in the scanning electron microscope to determine the response of the wear surface to electrotribological loading. Results and Discussion A. Binderless Copper-Graphite Composite Brush Material Preliminary results of the tribological characterization of the binderless copper and graphite powder-based composites were previously published by Owen et al (18), Wang et al (19), and by Eliezer et al (20, 23). The lox\ speed pin on disk tests (0.6 m/s) of the copper-graphite composite on 4340 steel showed frictioi. coefficients in the range 0.133 to 0.207. At high speeds the wear of the binderless composite is significantly lower than the commercial composite material now being used. At 200 m/s values of 0.2-0.3 x 103 mm3/Mm versus 2 x 103 mm 3 /Mm, are obtained over 6 km slieing distances. The peak interface temperatures observed by video thermography during these tests were 450 'C to 620 'C in the downforce range of 44.8 N to 134.4 N.
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Based upon previous work in our laboratory, Everett (24) reported on the pulse duty (
